Dense sic ceramic products

ABSTRACT

A process for the liquid phase sintering of silicon carbide, comprising forming a shaped, consolidated powder body which, not allowing for binder, comprises a powder mixture containing at least 75 wt % silicon carbide and from 1 to 25 wt % (calculated as Al 2  O 3 ) of a powder comprising a source of aluminum selected from alumina, precursors for alumina and mixtures thereof. The body is heated in a non-oxidising atmosphere to a sintering temperature of from 1500° C. to 2300° C. to form a liquid phase and a resultant liquid phase sintered body. In said heating step, the body is heated in the presence of a source of magnesium which is distinct from the source of aluminum and comprises at least one of magnesia, precursors for magnesia, magnesium vapour and combinations thereof, whereby said liquid phase produces secondary oxide constituent.

TECHNICAL FIELD

The invention relates to the production of dense articles of SiC.

BACKGROUND OF THE INVENTION

Shaped articles comprising polycrystalline SiC are known. They arecharacterized by excellent physical properties, such as high resistanceto thermal shock, abrasion and oxidation, together with high levels ofstrength and thermal conductivity. It is this combination of propertieswhich makes SiC materials leading candidates for engineeringapplications. However, this combination of properties only concurs inhigh density materials.

During high temperature heat treatments of prerequisite powder compacts,a reduction in the surface energy of the system can occur. The reductionin surface energy is through the diffusion of atoms by either grainboundary diffusion and subsequent densification, or by grain growththrough surface diffusion mechanisms with virtually no macroscopicdensification. At the high temperatures required for the sintering ofSiC powder compacts, surface diffusion typically prevails over grainboundary diffusion. This results in coarsening of the SiC grains in apowder compact with little macroscopic densification taking place.

The oldest process for production of dense articles of SiC is that ofreaction sintering, in which silicon liquid or vapour is infiltratedinto a compacted body of SiC powder and C. The Si reacts with the C toform SiC in situ which bonds the powder particles. However, this processtypically leaves from 8 to 12 volume percent of free Si, which sets amaximum operating temperature of about 1300° to 1400° C. for theresultant densified article.

In more recent times, attention has been directed to the use of certainadditives which promote grain boundary diffusion over surface diffusionfor pressureless sintering of SiC. However, apart from B or certaincompounds thereof, found to be effective in increasing grain boundarydiffusion, there does not appear to have been any successful proposal,at least in terms of commercial utility. Moreover, even with use of B ora B compound, problems still exist.

In use of B or a B compound, C usually is added as disclosed in U.S.Pat. Nos. 4,004,934, 4,041,117 and 4,108,929 all to Prochazka and U.S.Pat. No. 4,124,667 to Coppola et al. It is indicated that the C reducesthe surface SiO₂ layer on the SiC powder to SiC and CO. In U.S. Pat. No.4,041,117, Prochazka suggests that the SiO₂ can halt densification ofSiC compacts so that little or no shrinkage can occur. Prochazka alsosuggests that the addition of C can limit exaggerated grain growthduring densification. However, he further indicates that grain growthcan only be inhibited by strict control of temperature and pressurewithin narrow limits. Also, the final product usually contains Cparticles in the microstructure which can lead to degradation ofmechanical properties of the product.

The literature on effective sintering aids for SiC powder, other than Bor B compounds, is credited as having commenced with Alliegro et al, J.Amer. Ceram. Soc. 39 [11] 386-89 (1056). This reference discloses that1% Al addition to α- or β-SiC powder enables densification byhot-pressing to about 98% of the theoretical density. The β-SiC powderwas synthesised from a Si/C mixture, in which case, the Al usually wasadded to that mixture as oxide that was reduced during the synthesis.With use of α-SiC powder, the Al evidently was added as the metalpowder. Alliegro et al report that Fe, Li, Ca and Cr also aideddensification, but that Mg, Ta, Co, Ba, Mo, W, Sr and Cu were notbeneficial whether used alone or with Al.

Artemova et al, in Neoroanicheskie Materialy, Vol. 10, No. 12, pp2228-9, Dec. 1974, report on the preparation of a densified product byshock compression employing an explosive charge. Powdered SiC/Al₂ O₃mixtures ranging from 10/90 to 90/10 mole percent, in 10 mole percentincrements, were used and attained in excess of 98% of the theoreticaldensity for the mixtures. This mode of densification, having somesimilarity to hot-pressing, suggests the suitability of Al₂ O₃ as anadditive in SiC densification by more conventional procedures. However,Artemova et al report that it has not yet proved possible to densifymixtures of SiC/Al₂ O₃ at all by the usual methods.

Lange, J. Mater. Sci. 10 [1975] 314-320, reports on the hot pressing ofboth α- and β-SiC powder with use of Al₂ O₃ as a densification aid.While only quite small additions of Al₂ O₃ were used, ranging from 0.01to 0.15 volume fractions, densities up to and greater than 99% of thetheoretical density were achieved. Densification was attributed to aliquid phase which formed at high temperatures. However, the use of Al₂O₃, even at such low levels, was reported to result in large secondphase streaks of Al₂ O₃ of up to several millimeters long and usually 10to 15 μm wide. For brittle materials such as ceramics, the presence of aflaw, such as a crack, pore or inclusion can result in stressconcentration leading to failure. Streaks as reported by Lange wouldinevitably be deterimental to the physical properties of the densifiedSiC, as they greatly increase the defect size in the material.

It was speculated by Lange that the streaks of Al₂ O₃ were the result oflaminar void spaces present in the cold pressed specimens. Possiblesolutions to eliminate or reduce the occurrence of the streaks was toemploy a "sandwich" approach using layers of thinner bodies which, whencompacted, formed thicker bodies. This technique would be limited toprocedures such as hot pressing. Another technique proposed was graingrowth of the SiC grains. Under industrial conditions, the presence ofsuch voids is always possible with the probability of their occurrenceincreasing with increasing thickness of the component. Grain growth ofSiC to aid the removal of such streaks may prove difficult to control inpractice. Excessive grain growth is a problem associated with some ofthe techniques employed for pressureless sintering of SiC. This isconsidered to be a disadvantage in the use of Al₂ O₃ as a densificationaid. No indication was given by Lange as to whether Al₂ O₃ would be aneffective aid for the pressureless sintering of SiC.

Omori et al, J. Am. Ceram. Soc. 65 [1982] C-92, disclose the use ofoxide additives, viz. Al₂ O₃ and Y₂ O₃, in the Pressureless sintering ofβ-SiC powder. The oxides were used separately at 10 wt %, and incombination to a total of 10 wt % at ratios of 4:1, 3:2, 1:1, 2:3 and1:4. Densification was achieved with 10% Al₂ O₃, but only with 4%shrinkage and a relative apparent density of about 75%. With decreasingAl₂ O₃ content, densification was enhanced to about 97% of thetheoretical density at an oxide ratio of 1:1, but the level of the firedbulk density thereafter decreased and, with 10 wt % of Y₂ O₃ alone, nobenefit was obtained over β-SiC alone. Omori et al reasonably concludethat Al₂ O₃ enhances pressureless sintering despite its partial lossattributed to sublimation, but that Y₂ O₃ does not. However, the resultsdo suggest that, to a degree, Y₂ O₃ improves the beneficial effect ofAl₂ O₃. Omori et al report the loss of Al₂ O₃ on sintering, with aresidue of this oxide being determined by chemical analysis but not byX-ray diffraction.

A more recent study by Negita, J. Am. Ceram. Soc. 69 [12] C-308-C-310(1986), reports on the selection of suitable densification aids for thesintering of SiC. Using thermodynamic arguments, Negita reported thatmetal additives such as B, Al, Fe, Ni and Mg should be effectivesintering aids for SiC and that this had been found to be the caseexperimentally. In relation to B, Al and Fe, this accords with the workof others, as detailed above. On the basis of the same arguments, Negitareports that metal oxides, including Al₂ O₃, BeO, Y₂ O₃ HfO and rareearth oxides, should be effective densification aids, and that this wasborne out experimentally at least for Al₂ O₃, BeO, Y₂ O₃, La₂ O₃, Ce₂ O₃and ThO₂.

In contrast to the oxides listed in the previous paragraph, Negitareports that metal oxides including CaO, MgO and ZrO₂ are indicated notsuitable as they tend to decompose SiC. In addition, Negita suggeststhat the use of C with metal oxides is indicated as beneficial in thecase of Al₂ O₃, BeO, Y₂ O₃, CaO, ZrO₂, HfO₂ and rare earth oxides.

The use of Al₂ O₃ as a densification aid in the pressureless sinteringof SiC powder is disclosed in U.S. Pat. No. 4,354,991 to Suzuki et al.The proposal of this reference is to mould a mixture of anoxygen-containing Al-compound, which can be converted into Al oxide byheating in a non-oxidative atmosphere at a ratio of 0.5 to 35 wt % Al₂O₃, with the remaining ceramic material substantially being SiC. Suchmoulded mixture is subjected to pressureless sintering in anon-oxidative atmosphere at 1900° C. to 2300° C. Despite the requirementthat the oxygen-containing compound is one which can be converted intoAl oxide, it evidently is envisaged that the compound can be Al oxide.However, a number of disadvantages, of which some are confirmed by ourwork on the pressureless sintering of mixtures of SiC and Al₂ O₃, areapparent from U.S. Pat. No. 4,354,991.

The fired bulk densities obtained by the teaching of U.S. Pat. No.4,354,991 are relatively low, and also subject to substantial variationwith firing conditions. Also, the sintering times are relatively long,ranging from a preferred minimum of 2 hours up to 24 hours, with 3 to 5hours being typical even with relatively small samples. In a continuousprocess for densification of SiC powder, such reaction times wouldresult in lower production rates. Furthermore, another problem exists inthe preferment for control and maintenance of Al species in the firingfurnace atmosphere for long periods of time required for sintering.

No mention is made in U.S. Pat. No. 4,354,991 of the formation ofstreaks of Al₂ O₃ as reported in the above-mentioned article by Lange,even though such defects are likely to be a characteristic of the use ofAl₂ O₃ alone. As suggested by Lange, long soak times may be required toeliminate such streaks, and this possibly explains the relative longsintering times taught by U.S. Pat. No. 4,354,991. However, as indicatedherein, the use of such sintering times is disadvantageous.

We have found that a further apparent characteristic of the use of Al₂O₃ alone as a sintering aid for SiC powder is the tendency for zoning,particularly in the production of relatively large articles. That is, wehave found that use of Al₂ O₃ alone has a pronounced tendency to producea well densified outer layer enclosing an internal core which canexhibit substantially less densification. Where zoning occurs, thearticle is at least less than optimum. Also, internal stress due to thezoning can result in the article exhibiting cracks or, in extreme cases,the article can fail completely with the outer layer spalling.

The tendency for zoning with the use of Al₂ O₃ alone as a sintering aidfor SiC powder, as taught by U.S. Pat. No. 4,354,911, is believed to bedue to the difficulty of producing a sufficient volume percent of liquidphase at an appropriate temperature. This difficulty may also explainthe tendency for streak formation as reported by Lange, or streakformation may exacerbate the difficulty in achieving a sufficient volumeof liquid phase. As is known, efficient liquid phase sintering requiresnot only the formation of a liquid phase at a suitable temperature, butalso the presence of that phase in a sufficient volume over a suitabletemperature range.

In the proposal of U.S. Pat. No. 4,354,911, formation of a suitableliquid phase is not possible simply by melting of Al₂ O₃, except atextremely high temperatures. The melting point of Al₂ O₃ is about 2015°C., while decomposition and loss by volatilization of decompositionproducts thereof commences below that temperature, as recognised bySuzuki et al and also taught by the above-mentioned article by Omori etal. Despite the sole addition of Al₂ O₃ as a sintering aid, SiO₂ also ispresent as an impurity layer up to about 2 wt % on finely divided SiCpowder (unless previously removed), and the SiO₂ can facilitate theformation of a liquid phase at a temperature below the melting point ofAl₂ O₃.

Reference to the phase diagram for the SiO₂ -Al₂ O₃ binary system showsa eutectic composition at about 93% SiO₂ --7% Al₂ O₃ which has a meltingpoint at about 1595° C. Thus, assuming that the rate of heating to thesintering temperature range of 1900° to 2300° C. is not excessive,solid-solid diffusion between the separate Al₂ O₃ and SiO₂ can give riseto an initial small volume of liquid at temperatures above 1595° C.Also, SiO₂ melts at about 1730° C. and, assuming that the SiO₂ is notpreviously volatilized or decomposed, as tends to occur, a small volumeof SiO₂ -containing liquid phase can be formed above 1730° C. and thiscan increase in volume by taking up Al₂ O₃ by liquid-solid diffusion.

In the method taught by Suzuki et al in U.S. Pat. No. 4,354,911, thelower level of Al₂ O₃ addition is 0.5 wt %, corresponding to an SiO₂ toAl₂ O₃ ratio on the Al₂ O₃ rich side of the eutectic of the SiO₂ -Al₂ O₃binary system. That is, when allowance is made for 2.0 wt % SiO₂ beinghigh and 0.5 wt % Al₂ O₃ being a minimum, it is apparent that a bestpossible ratio is about 80% SiO₂ : 20% Al₂ O₃. A lower SiO₂ content or ahigher Al₂ O₃ content rapidly advances that ratio away from the eutecticcomposition to increasily richer Al₂ O₃ contents. At only 2.0 wt % Al₂O₃, the ratio is at least at the mid-point of the SiO₂ -Al₂ O₃ phasediagram. At 4.0 wt % Al₂ O₃, the ratio is such that little, if any,liquid previously formed will remain, with further liquid then not beingformed until a temperature of about 1840° C. is achieved. That is, withan Al₂ O₃ content of at least 4.0 wt % Al₂ O₃, any liquid initiallyformed will be substantially lost, due to precipitation of a corundum ormullite solid phase having a melting point of about 1840° C. However,given that an Al₂ O₃ addition of only 0.5 wt % still is on the Al₂ O₃-rich side of the eutectic composition, at least a proportion of anyinitially formed liquid with less than 4.0% Al₂ O₃ additions willsimilarly be lost due to precipitation of corundum or mullite. Theseproblems are further exacerbated by the tendency for SiO₂ and Al₂ O₃ todecompose and to be lost by volitization of their decomposition productsat temperatures approaching 1840° C., making it very difficult toproduce, or produce and retain, a significant volume of a liquid phase.Also, Al₂ O₃ present at a level significantly in excess of 4.0 wt % willnot be able to be taken fully into solution below at least about 1840°C., with the temperature at which this is possible rapidly increasingwith the level of Al₂ O₃ addition to about 2015° C. Moreover, if thereis only alumina present, a liquid phase cannot be formed below themelting point of Al₂ O₃, that is, below about 2015° C., and even then, aliquid will only form if some Al₂ O₃ is retained until that temperatureis attained.

The precipitation of corundum or mullite from initially formed liquidmay explain the streaks of Al₂ O₃ reported in the above-mentionedarticle by Lange. The streaks are referred to by Lange as suggesting a"frozen liquid". This may well have resulted from corundum or mulliteprecipitated from an initially formed liquid, and only partiallyremelted on heating at about 1840° C.

For the temperature range of 1900° C. to 2300° C. taught by Suzuki et alin U.S. Pat. No. 4,354,911, and the addition of Al₂ O₃ alone at from 0.5to 35% as a sintering aid for SiC, it therefore is extremely difficultto achieve a liquid phase at all, let alone one in a sufficient volumefor efficient liquid phase densification. As the Al₂ O₃ level increasesabove 0.5%, the temperature at which fully liquid SiO₂ and Al₂ O₃ ispresent also increases, and the volume of liquid able to be producedbelow the 1840° C. solidus decreases. Particularly above about 4% Al₂O₃, it can be necessary to use a temperature substantially above 1900°C. in order to achieve any significant volume of liquid at all.

A further disadvantage of the proposal of U.S. Pat. No. 4,354,991 arisesfrom the strong preferment for use of β-SiC powder, rather than, α-SiCpowder. β-SiC is not as readily available as α-SiC as produced by theconventional Acheson process for the manufacture of SiC grit. Thatprocess accounts for a major portion of world-wide production of SiC andα-SiC is readily available and is a commodity traded on the worldmarket.

In International patent specification PCT/US88/00040 (W088/05032),Fuentes discloses the pressureless sintering of SiC powder, using as asintering aid a combination of Al₂ O₃ and CaO. Fuentes recognises thatwith use of Al₂ O₃ alone as a sintering aid for SiC, the liquid phasenecessary for sintering is deficient in volume and/or forms too slowly.He therefore teaches use of a sintering aid mixture which produces aliquid phase at from 1815° to 1855° C. and comprises Al₄ O₄ C and Al₂OC. However this liquid phase, which also can be generated by use of Al₄O₄ C and Al₂ OC ab initio, itself forms at an excessively hightemperature for optimum densification. In addition, as reported byFoster et al, J. Am. Ceram. Soc. 39 [1956] 1-11, Al₄ O₄ C, Al₂ OC andAl₄ C₃ are very unstable towards both moisture and oxygen. The presenceof these species in the product resulting from the process taught byFuentes is very undesirable, and to be expected to greatly degrade theperformance and severely limit the utility of the product. The processand product as disclosed by Fuentes therefore presents significantdisadvantages.

In contrast to the prior art discussed above relating to thedensification of SiC to produce bodies of high density approaching thetheoretical density, the use of oxides for the bonding of SiC grits toform refractory bodies also has been considered. Thus, in U.S. Pat. No.2,040,236 to Benner et al, the use of a bonding material of Al₂ O₃together with either CaO, MgO or a mixture of CaO and MgO was consideredfor use in bonding SiC grit in producing a refractory body. Benner et alteach the heating in a non-oxidizing atmosphere of a pressed mixture ofSiC grit and such bonding material. The heating was to a relatively hightemperature, at which the bonding material softened to undergo incipientfusion. However, the rate of heating to temperature was rapid, such asabout 35 minutes. Also, it is emphasised that the time at temperaturewas to be short so that, while sufficient to soften the bondingmaterial, recrystallization of SiC could be avoided. Furthermore, thetime at temperature was to be short so that the bonding material did noteither decompose or react with the SiC.

The suitable SiC grit proposed by Benner et al ranged from 14 mesh toless than 80 mesh, but with coarse, medium and fine size fractions. Thus40% was--40 mesh+36 mesh (ranging from less than about 1170 μm to about410 μm); 10% was--40+70 mesh (ranging from less than about 370 μm toabout 190 μm); and 50% was of--80 mesh (ranging down from about 180 μm).While only the small sub-micron portion of the fine size fraction wouldbe appropriate for densification as required by the prior art discussedabove, Benner et al report production of a useful refractory comparedwith use in a similar context of other bonding materials. Theirrefractory is said to have been very dense and of lower permeability inthat context. Microscopic examination (as applicable in 1932) is statedto have shown the product to exhibit pores only partially filled withbond material, while the refractory was permeable to gases. In thisregard, the disclosure of Benner et al is devoid of any indication thatmacroscopic densification of the body occured. Also, the SiC particlesof the grit, as confirmed by reference to it as a filler, in essence wasbonded in a matrix of the bonding material, with the latter evidentlyremaining in essentially the proportion of, for example, 5 to 10% inwhich it was added to the mixture.

The teaching of Benner et al detailed in the preceding two paragraphs isappropriate for the bonding of SiC grit, but does not provide guidancerelevant to liquid phase densification SiC powders. That is, they areseeking to produce refractories by bonding SiC grit particles in amatrix. The matrix acts in effect as a cement or glue (in the generalsense of these terms) which encapsulates and isolates the SiC gritparticles without decomposing or reacting with the SiC of theseparticles. In contrast, liquid phase sintering necessitates finer SiCpowder of a compact being densified being taken into solution andsubsequently precipitated, such as onto larger SiC grains, with the endproduct having clearly defined grain boundaries between SiC grains andany second phase. In effect, Benner et al teach use of a passive bondingmaterial which softens to form a matrix, whereas liquid phase sinteringrequires the presence of an active liquid phase which is formed by theassistance of sintering aids.

The non-oxidizing atmosphere proposed by Benner et al was required toinert to both the SiC and the bond material. Carbon monoxide isindicated as being satisfactory relative to Al₂ O₃, MgO and CaO andtheir mixtures. However, where SiO₂ was a principal constituent of thebond material, a more inert atmosphere such as nitrogen or helium waspreferred.

Further, in U.S. Pat. No. 4,829,027 Cutler et al disclose liquid phasesintering of SiC with use of a rare earth oxide and Al₂ O₃ ; the rareearth oxide principally exemplified being Y₂ O₃ as in the Omori et alreference considered above. The disclosure of this reference emphasisesthe importance of attaining a liquid phase at a relatively lowtemperature, in achieving densification by pressureless liquid phasesintering of SiC, substantiating our findings in relation to adissimilar system based on use of Al₂ O₃.

Finally, Japanese patent application 01230472, public disclosure No.89-230472, by Kurosaki Refractories Co. Ltd., proposes the production ofSiC sintered products using alumina/magnesia spinel (i.e. MgAl₂ O₄) as asintering aid. Kurosaki teaches that when spinel alone is used as asintering aid, magnesia will evaporate preferentially from the surfaceof the spinel powder grains, leaving grain surfaces covered with a layerof Al₂ O₃. During sintering, a liquid phase is said to form attemperatures of 1900° C. and above; this being seen as beneficial inresulting in little likelihood of deterioration of the excellent hightemperature characteristics inherent in SiC. In this regard, theteaching of Kurosaki is to avoid a liquid phase being formed at fairlylow temperatures, a matter on which they are at variance with the clearteaching of Fuentes, Cutler et al and our research.

A disadvantage of the teaching of the Kurosaki proposal is the relianceon relatively expensive spinel as the sintering aid, particularly as inexcess of 5 wt % spinel is necessary for optimum results. In thismatter, the same disadvantage exists with the proposal of Cutler et alin their reliance on expensive rare earth oxides. However, further majordisadvantages exist with the proposal of Kurosaki. The first is thatarising from the loss of MgO to which they refer since, with increasinglevel of spinel, the resultant weight loss will be increased; with apossible maximum of about 9.9 wt % due to this factor alone at 35 wt %spinel. However, as made clear by the work of others considered above,and also substantiated by our findings, these weight losses are likelyto be exacerbated by additional loss of SiO₂. Al₂ O₃ and SiC. A furtherimportant disadvantage is that, due to the spinel grains becoming coatedwith Al₂ O₃, any liquid phase initially tending to form will requireslow solid-solid diffusion, followed by dissolution of Al₂ O₃ andspinel, with this occurring to any significant extent in a reasonabletime only at temperatures substantially above 1900° C. This will lead toessentially the same problems in achieving a sufficient volume of liquidphase necessary for efficient liquid phase densification, as discussedabove in relation to the teaching of Suzuki et al.

SUMMARY OF THE INVENTION

The present invention seeks to provide an improved form of densearticles produced from SiC powder, and to an improved method ofproducing such articles. In particular, the present invention isdirected to providing such articles utilising Al₂ O₃. However, in thepresent invention, the use of Al₂ O₃ is under conditions which overcomeproblems, exemplified by the prior art and confirmed by our findings,encountered with the use of Al₂ O₃ alone.

As detailed above, use of Al₂ O₃ alone as an additive in densificationof SiC powder necessitates use of relatively high sintering temperaturesand relatively long sintering times. However, even with recourse to suchconditions, we have found that it can be difficult to achieve asatisfactory product. Indeed, unless other conditions are satisfied suchas use of a powder bed or coating as taught by Suzuki et al in U.S. Pat.No. 4,354,911, the resultant product can exhibit minimal, if any,densification and low strength, such that the product can readilycrumble. Also, even where a useful level of densification is achieved,this can be limited to an external surface layer, with the interior ofthe product being less satisfactorily densified and the sectionedproduct exhibiting a macroscopically visible cored structure. Theinterior of a product exhibiting such structure can have a relativelyhigh degree of densification, even comparable to that of the surfacelayer. However, we have found that densification and composition aldifferences, or both, between the surface layer and core can result inthe product exhibiting cracks which reduce the mechanical properties ofthe product. These differences can be such that the product as formedhas failed, for example by propagation of cracks or spalling of thesurface layer from the core, due to stress generated in the product oncooling from the densification temperature. Moreover, streaks of Al₂ O₃,such as reported in the above mentioned article by Lange, can be presentin the microstructure of the product, and it is believed that suchstreaks can facilitate crack formation or propagation. The disadvantagesdue to the use of Al₂ O₃ alone as an additive for densification of SiCpowder can be ameliorated, at least to a degree, by use of a relativelyhigh sintering temperature and, more importantly, use of a relativelylong time at the sintering temperature of at least about 2 hours, buttypically at least 3 hours. However, such expedients substantiallydecrease production rates, thereby increasing the cost of production.They also result in an increased loss of Al₂ O₃ by decomposition, andsubstantially increase the requirements for control and maintenance ofspecies, in the atmosphere in which sintering is conducted, intended toprevent or offset such loss.

We have found that the above problems encountered with use of Al₂ O₃alone as an additive for densification of SiC powder can be overcome byuse of at least one of Al₂ O₃ and a precursor for Al₂ O₃ (hereincollectively referred to as the Al source) in combination with at leastone of MgO and a precursor for MgO (herein collectively referred to asthe Mg source). However, at the outset, it is to be understood that thepresent invention is concerned with a Mg source which is distinct fromthe Al source, rather than one which, as in the teaching of Kurosaki, isable to be recognised as an intimate constituent of a material such asspinel.

It also is to be understood that there are important variants of theinvention. In a first variant, a combination of the Al and Mg sources isprovided as an additive in a powder mixture with SiC powder as preparedto form a powder compact for sintering. In a second variant, only partof the Al source requirement is provided in the powder mixture to formthe compact, with the balance of the required Al source being formed inthe compact during heating to the sintering temperature from Al-speciesprovided in the atmosphere in which the heating is conducted. In a thirdvariant, not more than part of the Mg source requirement is provided inthe powder mixture for the compact, with the balance or all of therequired Mg source being similarly formed in the compact during heatingfrom the Mg-species provided in that atmosphere. The second and thirdvariants can be used in combination, with part of the Al sourcerequirement and part or all of the Mg source requirement being formed inthe compact from Al-species and Mg-species provided in the atmosphere.Also, provision of Al-species and/or Mg-species in the atmosphere can beof benefit during the first variant, as the species act to offset lossof Al and/or Mg source from the compact, at least at highertemperatures.

The Al-species and Mg-species, able to be provided in the atmosphere inwhich the compact is heated, comprise species able to be present in theatmosphere in a gaseous condition. However, the species also need to besuch that they will permeate the compact and react with a constituent ofthe compact to form Al₂ O₃ and MgO. In general, the species compriseelemental AL, Al₂ O and elemental Mg. The species may be generated inthe furnace in which the compact is heated, as hereinafter detailed, orthey can be charged to the furnace in gaseous form from a suitablesource, or respective source, external to the furnace.

The beneficial effects of use of an Al source in combination with an Mgsource, as set out in more detail in the following, is surprising inview of the prior art considered above. Thus, while the above-mentionedarticle by Negita reports on the utility of Al₂ O₃ as a suitableadditive for densification of SiC powder, MgO is reported as not being asuitable additive. Moreover, Negita reports in the above-identifiedpaper that CaO which, as will be appreciated, is chemically equivalentto MgO in most contexts, also is not suitable. Fuentes, in theabove-mentioned International patent specification PCT/US88/00040(W088/05032) reports that use of Al₂ O₃ and CaO can be used incombination as an additive in the pressureless sintering of SiC, butthat this combination results in an undesirable secondary phase of Al₄O₄ C and Al₂ OC. We have found that use of Al₂ O₃ and MgO in combinationas an additive in pressureless sintering of SiC does not give rise tosuch oxycarbide secondary phase while it also gives rise to importantdifferences of practical benefit not available with the process ofKurosaki.

Moreover, while Benner et al in U.S. Pat. No. 2,040,236 teach that acombination of Al₂ O₃ and MgO has utility in bonding SiC grit, theirteaching indicates strongly that such combination would be unsuitablefor densification of SiC powder by pressureless sintering. Thus, whileit would be expected that the temperatures taught by Benner et al forbonding SiC grit are at a level appropriate for such sintering of SiCpowder, the rapid rate of heating to, and short time at, suchtemperatures are not suitable for sintering of SiC powder. Theirteachings on the avoidance of conditions which result in substantialrecrystallization of SiC, in reaction of the Al₂ O₃, MgO or both withSiC and in substantial decomposition of the Al₂ O₃ and/or MgO areindicative of conditions which are suitable merely for cementing SiCgrit particles in a matrix of bonding material and which areinconsistent with requirements for sintering SiC powder. The teaching ofBenner et al is in relation to SiC grit, and problems ofrecrystallization of SiC and its reaction with such mixture would beexpected to be substantially more severe with sintering of SiC powder,given that SiC powder will be very much finer and of substantiallygreater surface area. In stark contrast, the clear indications are thatin the present invention densification, which is substantial, isfacilitated by the mechanism of liquid phase sintering. In thismechanism, an important feature is the dissolution and re-precipitation(recrystallization) of finer SiC particles, which is directly opposed tothe teaching of Benner et al. The SiC powder required for sinteringtypically is of a particle size less than 10 μm, such as of sub-micronsize on average. That is, the very much finer particle size and, hence,very much larger surface area, of SiC powder for sintering, comparedwith the grit of Benner et al, would be expected to result in excessiverecrystallization of SiC and loss of SiC by reaction, leading todegradation of physical properties even if macroscopic densification wasfound to result. Also, under conditions for pressureless sintering ofSiC powder with use of Al₂ O₃ and MgO in combination as an additive, theteaching of Benner et al is that the additive would be lost bydecomposition.

DETAILED DESCRIPTION OF THE INVENTION

The present invention provides a process for the production of a denseSiC product, comprising forming a consolidated powder compact bypressing a powder mixture containing at least 65 wt % SiC powder and atleast 1 wt % of a powder comprising an Al source, and heating thecompact in a non-oxidising atmosphere to a sintering temperature of from1500° to 2300° C., said compact being heated in the presence of at leastone of an Mg source and Mg-species, with the time at said sinteringtemperature being sufficient to achieve a required level ofdensification by liquid phase sintering.

Thus, according to the invention, there is provided a process for theliquid phase sintering of silicon carbide, comprising the steps of:

forming a shaped, consolidated powder body which, not allowing forbinder, comprises a powder mixture containing at least 70 wt % siliconcarbide and from 1 to 25 wt % (calculated as Al₂ O₃) of a powdercomprising a source of aluminium selected from alumina, precursors foralumina and mixtures thereof; and

heating the body in a non-oxidising atmosphere to a sinteringtemperature of from 1500° C. to 2300° C. to form a liquid phase and aresultant liquid phase sintered body;

the body, in said heating step, being heated in the presence of a sourceof magnesium which is distinct from the source of aluminium andcomprises at least one of magnesia, precursors for magnesia, magnesiumvapour and combinations thereof, whereby said liquid phase producessecondary oxide constituent.

The process of the invention enables production of a sintered bodyhaving a good degree of uniformity of physical properties. Also, thebody typically exhibits a microstructure in which SiC grains aresubstantially equi-axed and of greater grain size uniformity than theSiC powder used. The microstructure shows a reduction in the proportionof fine SiC grains relative to the powder, consistent with the fines ofthe powder and sharp edges of larger particles having been dissolved.Grains of SiC in the microstructure, in addition to being equi-axed,show a degree of rounding consistent with precipitation of dissolved SiCon and between larger grains, and formation of neck regions at SiC grainboundaries. In these regards, the aspect of the SiC grains is understoodto be quite distinct from the fine structure of interlaced tabular, orcrossed plate, crystals obtained by the teaching of Kurosaki.

Depending on the level of secondary oxide constituent in the sinteredbody, the microstructure of the invention can be of duplex form. Thatis, the secondary oxide can be present as an inter-connected networkthroughout sintered grains of SiC, with a substantial proportion ofadjacent SiC grains exhibiting well defined SiC to SiC grain boundaries.The secondary oxide is rich in Al, and can substantially comprise Al₂O₃. The secondary oxide can, and typically does contain Mg. However,while the ratio of Mg to Al in the compact as formed can be as high as1:2, the secondary oxide typically has a lower ratio (i.e. richer in Al)such as at least 1:3. Also, even though the Mg to Al ratio in thecompact can be lower than 1:2 (richer in Al) and typically is, the Mg toAl ratio in the resultant product usually is lower still.

Our research has found that the secondary Al rich oxide constituentproduced by in situ reaction of stoichiometric additions of Al sourceand Mg source does not necessarily result in useful bodies especiallyfor articles with thicker cross sections. Indeed, it has been shown thatuseful bodies are produced by obtaining a desired Mg to Al ratio in thefired body which is less than 1:3 at low Al source contents and 1:6 atthe upper levels of Al source addition. The present process hassubstantial flexibility in the allowable ratios of Al source and Mgsource in the compact as formed. There is no requirement for the loss ofMg at elevated temperatures to allow the formation of a liquid phase bythe reaction of the Al₂ O₃ formed as a decomposition of spinel, as inKurosaki, with spinel then to initiate the densification of SiC. It hasalso been found that the use of powder beds can result in an increase inthe amount of Al in the final body. This allows the densification ofbodies initially low in Al to proceed as a result of transfer of Al tothe body, changing the Mg to Al ratio to a more favourable level suchthat densification can proceed. From this it is found that when thelower limit of Al source is used, sufficient pickup and diffusion acrossthe body occurs to enable uniform densification to proceed. For thickerbodies it is not always possible and practical to allow adequate time toeffect such diffusion and sufficient pickup. The end result is theproduction of bodies with porous cores. This may be useful in its ownright but this is not always the case and hence an upper limit on thethickness can exist to allow the production of high density bodies. Inthe process of Kurosaki, the liquid phase is reported to only appear atelevated temperatures. The SiO₂ inevitably present on the SiC as animpurity is most likely to be essentially lost from the system and noteffectively utilised. In contrast, the present process, demonstrates theadvantages of low temperature holds which allow the reaction of MgO andAl₂ O₃ with SiO₂ present. As the temperature is increased to the middletemperature ranges considered, this phase will provide liquid inincreasing amounts to effect densification by well known liquid phasesintering processes which involve particle rearrange of SiC grains inthe presence of the liquid and secondly by the solution precipitation ofSiC. In addition, research has shown that it can be difficult to effectloss of Mg from the compacts during the densification operationespecially for samples with thicker cross section or when furnaceloadings are high. It has also been found that the decomposition anddeposition of Mg in the system can interfere with the operation of thefurnace. These factors highlight the problems inherently associated withthe use of the spinel as proposed by Kurosaki and the advantages of theprocesses as disclosed herein.

The SiC powder preferably is substantially free of SiC particles largerthan 10 μm and has an average particle size substantially less than 10μm. Most preferably, the SiC powder has an average particle size lessthan 2 μm. The SiC powder may comprise α-SiC of any polytype, β-SiC,amorphous SiC, or mixtures thereof.

At least part, preferably a major part, of the required level of the Alsource is present in the consolidated body (hereinafter referred to as a"compact") as formed. The Al source in the compact as formed mostpreferably comprises Al₂ O₃, in any of its available forms including α-and γ-alumina, although a precursor which generates Al₂ O₃ on heatingcan be used instead. Suitable precursors for Al₂ O₃ include Al(OH)₃,Al(NO₃)₃, 3Al₂ O₃.2SiO₂, Al₂ O₃.SiO₂, AlO(OH), organo-metallic salts ofAl including fatty acid salts, other Al compounds which decompose onheating to yield Al₂ O₃, and mixtures of these compounds with or withoutAl₂ O₃.

While it is required that the compact as formed contains at least partof the required level of Al source, it is not necessary that it alsocontains at least part of the required Mg source. It is preferred thatthe compact as formed does contain at least part of the required levelof Mg source and, to the extent that this is the case, the Mg sourcemost preferably comprises MgO. However, Mg source in the compact mayconsist of or include a precursor which generates MgO on heating.Suitable precursors for MgO include MgCO₃, Mg(OH)₂, Mg(NO₃)₂,organometallic salts of Mg including fatty acid salts, and mixtures ofthese compounds with or without MgO.

Al source and Mg source present in the compact as formed preferably areless than 10 μm where comprising a powder. However, at least part of theAl source, the Mg source, or both, need not comprise a powder, such aswhere comprising an organic precursor material. An Mg source comprisingorganic material can be beneficial, particularly in the case of a fattyacid salt such as a stearate, able to act as a lubricant in forming thecompact. Such organic material can be present in the compact as a filmcovering powder particles of the compact, rather than as a powder perse.

Part of the required level of the Al source, preferably a minor part, oradditional Al source, may be present in the furnace in which the compactis heated. Part or all of the required level of Mg source, or additionalMg source, may be present in the furnace. These conditions can apply asalternatives, or in combination. Al source, Mg source or both soprovided in the furnace may be in or comprise a particulate bed on or inwhich the compact is heated, or it may be in or comprise a coatingformed on or around the compact. Al source, Mg source or both present ina bed or coating can be of a powder subject to the same particle sizeconstraints as the SiC. However, such source of the bed or coating canmore readily be provided in a greater quantity than is permitted for itspresence in the compact, and source of larger particle size, as ispreferred, such as in the form of grit of up to 2 mm can be used in thebed or coating.

Whether or not substantially all of the Al and Mg sources are providedin the compact, sintering of the compact most preferably is conducted inthe presence of at least one of Al-species and Mg-species in thesintering atmosphere. Such species may be generated by passing asuitable atmosphere containing the species through the furnace from anexternal supply, by generating the species in the sintering furnace byheating a suitable bed or coating as described above, or by a suitableratio of compact mass to furnace volume. Use of a bed or coating ispreferred, with use of a bed being most preferred. Where a coating isused, it may comprise a coating formed on the compact or on a vessel inwhich the compact is positioned. The coating most preferably is formedfrom a slurry containing suitable Al source, Mg source or both, afterwhich a layer of the source is dried to form the coating by lowtemperature heating.

As made clear in the foregoing, at least part of the Mg source can bepresent in the compact as formed, and this is preferred. Indeed, it ispreferred that at least a major part of the Mg source requirement ispresent in the compact as formed. Where the compact contains all of theAl source and the Mg source, the compact is formed from a powder mixturecontaining up to 30 wt % of Al source plus Mg source (calculatedrespectively as Al₂ O₃ and MgO), with the balance substantiallycomprising SiC powder. The amount up to 30 wt % provided by the Alsource plus Mg source preferably has quantities of Mg source (as MgO)and Al source (as Al₂ O₃) such that the ratio of Mg to Al is in therange of 1:2 to 1:25, at levels of 1 wt % to 5 wt % Al source (as Al andin the range of 1:5 to 1:100 at levels of 5 wt % to 25 wt % Al source(as Al₂ O₃) for increasing contents of Al source in that range of 5 wt %to 25 wt %, the ratio of Mg to Al decreases in a substantially linearmanner from 1:2 at 5 wt % to 1:5 at 25 wt % Al source (as Al₂ O₃).

Where the Al source or the Mg source present in the compact is otherthan as Al₂ O₃ or MgO, respectively, it decomposes during heating to thesintering temperature to provide Al₂ O₃ or MgO and gaseous decompositionproducts. Such gaseous products readily are able to escape from thecompact, at least in the initial stage of the densification process, dueto its initial porosity. Where the Al source, the Mg source or eachsource is provided in a powder bed or in a coating, it decomposes duringheating to that temperature, initially to provide Al₂ O₃, MgO or bothwhere the source respectively is other than Al₂ O₃ or MgO and generatesvapour such as of Al₂ O, Mg or both in the atmosphere in which sinteringis conducted. It is found that the Al₂ O or Mg vapour readily is able topermeate the initially relatively porous compact by vapour diffusion,and to generate a suitable source of Al or Mg therein.

A principal role of Al source and Mg source provided in the furnace,such as in a bed or coating, is to generate Al-species and Mg-specieswhich permeate and are taken up in the compact to achieve a desiredlevel of Al₂ O₃ and MgO in the compact. However, in generating suchspecies, the source provided in the furnace serves a second useful rolein generating a substantial partial pressure in the furnace. Evidentlydue to more rapid heating of the source in the furnace relative toheating of the more dense compact, it appears that the source in thefurnace decomposes to provide the species in advance of substantialdecomposition of source powder in the compact. The resultant partialpressure of the species appears, as a consequence, to be able to atleast partially suppress decomposition of the corresponding source inthe compact. In relation to each of these roles, the same result can beachieved by charging to, and maintaining in the furnace, an atmospherefrom an external source and containing the required Al-species,Mg-species or both. The atmosphere is produced outside the furnace, suchas by heating the required Al source, Mg source or both, maintaining theatmosphere at a temperature sufficient to prevent condensation of thespecies, and passing the atmosphere through the furnace. Particularlywhere both Al-species and Mg-species are separately formed, the ratio atwhich they are present in the atmosphere, as passed to the furnace canbe selected. Also, depending on the furnace construction, it is possibleto monitor the ratio of the species discharging from the furnace, andthe ratio can be adjusted, if appropriate.

If the required Al source and Mg source are both provided in thecompact, it still is desirable but not necessary to utilise a bed orcoating or to generate Al- and Mg- species by charging these to thefurnace in an atmosphere generated externally of the furnace. In thiscase, Al₂ O₃ and MgO will tend to decompose and be lost from thecompact. This can be offset to a degree by an excess of Al source and Mgsource in the compact, to allow for the loss. However, as a morepractical alternative to use of a bed, coating or external supply ofatmosphere, it can be beneficial to ensure that the mass of compactsdensified in the furnace in a given firing is in a favourable ratio tothe volume of the furnace. That is, it is desirable that this ratio issuch as to ensure that Al- and Mg- species generated by the loss bydecomposition gives rise to a vapour pressure which restricts the lossto an acceptable level. In relation to the mass of compacts required, itwill be appreciated that this will depend on the mass to surface arearatio of the compacts, while the rate of heating and furnace design arefurther relevant factors.

While not wishing to be bound by a specific reaction mechanism, it isclearly indicated that MgO has an important role in achieving thebeneficial results provided by the process of the invention. Where theMg source is present in the compact as formed, MgO is present due to useof MgO as the source or is formed in situ by decomposition of aprecursor for MgO. In each case, our findings indicate that the MgOforms a transient liquid with SiO₂ and Al₂ O₃ at relatively lowtemperatures, such as with SiO₂ initially present due to surface layeroxidation of the SiC powder. Indications from our research are that suchtransient liquid forms at from 1300° to 1400° C. While initially presentin relatively small volume, the liquid is thought to result indissolution of other oxides present, such as Al₂ O₃, causing asubstantial increase in liquid volume. Also, while the liquid initiallyforms at a relatively low temperature, it appears to be stable and suchthat it is retained at higher temperatures at which a liquid necessarilyis to be present for efficient liquid phase densification.

Our research suggests that the transient liquid varies in compositionduring the course of heating to the sintering temperature, and withholding at that temperature. It seems clear that the liquid initially issubstantially of quasi-ternary composition comprising SiO₂, MgO and Al₂O₃. With increasing temperature above about 1300° C. the liquid takes upfurther MgO and Al₂ O₃. At still higher temperatures, with increasingtime or both, the SiO₂ evidently is progressively lost by decomposition,with at least some loss of MgO also being possible, resulting in an Al₂O₃ rich liquid which also can contain at least a residue of MgO. The endresult typically is a secondary oxide constituent rich in Al, such as γor α-Al₂ O₃. The secondary constituent produced in situ can, andtypically does, contain Al, but with an Mg to Al ratio of at least 1:3.However, the oxide constituent, despite resulting from the binary liquidphase, readily is able to comprise Al oxide with no detectable Mgcontent. Moreover, depending on the level of Al source and Mg sourceinitially present in compact as formed, and the extent of control overAl-species and Mg-species in the furnace atmosphere as detailed herein,the temperature and time of sintering can be such that substantially allMgO and thereafter substantially all Al₂ O₃ can be lost, resulting in asintered product containing little if any detectable oxide constituent.Despite this latter possibility, the process of the invention preferablyis conducted such that secondary oxide constituent, with or without adetectable level of Mg, is retained, as such constituent enhances thefracture toughness of the product. The loss of Al₂ O₃ can result in aminor quantity of Al metal being detected in the body. However, despiteloss of MgO and decomposition of Al₂ O₃, good densification still isachieved and this and the progressive change in the transient liquid tohigher Al to Mg ratios, indicates that the active Al- and Mg-speciesresponsible for densification are retained in the transient liquid phaseuntil densification is substantially complete or at least untildensification reaches a level at which another mechanism leads toattainment of a high density body.

Where the Mg source is at least partially present in a bed or coating,Mg vapour is generated and permeates through the compact by diffusion.Furthermore, it appears that oxygen-containing species also are presentin the furnace atmosphere. It is thought that the Mg vapour reacts withoxides contained in the compact, to generate MgO in situ, possibly withevolution of SiO. A similar mechanism is believed to occur with Mgvapour comprising Mg- species in an atmosphere passed to the furnacefrom an external supply. The MgO thus formed results in a low meltingpoint transient liquid with SiO₂ and Al₂ O₃, with that liquid thenresulting in ongoing dissolution of Al₂ O₃ and increasing in volume aspreviously explained in relation to the provision of the Mg source inthe compact as formed. In addition, it appears that part of the SiO₂initially present in the compact can decompose to form SiO gas and 1/2O₂, with either the oxygen reacting in the compact with the Mg speciesto form MgO or the SiO reacting in the compact with those species toform MgO in the presence of other oxygen containing species.

With provision of the Mg source either in the compact as formed, orderived from a powder bed, coating, or an external source of atmosphere,the resultant MgO in the compact is found to act as a flux. That is, theMgO substantially increases the fusibility of Al₂ O₃ and SiO₂ and, atlow temperatures, forms a transient liquid of sufficient volume which isretained at higher temperatures The SiO₂ is able to form a liquid phasebut, in the absence of MgO, SiO₂ tends to be lost by decomposition.

The invention also provides a sintered ceramic product, comprising asintered body having at least 65 wt % SiC with any secondary constituentcomprising not more than about 30 wt % secondary oxide constituent, andnot more than minor amounts of elemental Al, elemental Si and glassyphase. Any secondary oxide constituent present is rich in Al and maysubstantially comprise Al oxide. However, the constituent may includeMg, with an Mg to Al ratio of not more than 1:3. In contrast to theteaching of Kurosaki detailed above, any oxide constituent containing Mgis produced in situ by the use of Al₂ O₃ per se, or a precursortherefor, as distinct from spinel, such that the sintered body isproduced by liquid phase sintering resulting from a transient liquidphase produced at relatively low temperatures below that required forsintering.

The product of the invention preferably comprises at least 80 wt % SiCand most preferably at least 85 wt % SiC. However, the SiC content canexceed 98%. At SiC contents of at least 95 wt %, the product can appearto have a microstructure exhibiting only a single phase, orsubstantially only a single phase, with residual constituents other thanSiC evidently being in solid solution in that phase.

The product of the invention typically has a fired bulk density inexcess of 2.95 g.cc⁻¹. Fired bulk densities in excess of 3.00 g.cc⁻¹readily are able to be achieved, such as in excess of 3.10 g.cc⁻¹.Indeed, we have found that such densities of at least 3.15, such as upto and in excess of 3.25 g.cc are possible. Moreover, a good degree ofuniformity of fired bulk density, through a sintered body according tothe invention, is able to be achieved, while resultant low residualporosity can be substantially as required by control of the level ofsintering aid present in the initial compact as formed, the sinteringatmosphere and the sintering temperature and time.

The product of the invention also is able to exhibit a high level offracture toughness, as detailed herein. The fracture toughness increaseswith increasing level of secondary oxide constituent; frequently, butnot necessarily, with corresponding decrease in hardness. However, thehardness typically is in excess of about 18.5 GPa, and can range up toabout 26 GPa with significant levels of that constituent present. Thehigher levels of hardness are possible, particularly with decreasinglevels of secondary oxide constituent below about 5 wt % and with lowerlevels of residual porosity density. Also, as similarly detailed abovein relation to residual porosity, it is possible to achieve a requiredbalance between fracture toughness and hardness by control oversintering aid level, sintering atmosphere and sintering temperature andtime.

The product of the invention may be formed from a compact containingfrom about 1 to 25 wt % Al source (calculated as Al₂ O₃) However, it ispreferred that the Al source (as Al₂ O₃) is within the range of from 2.5to 20 wt %. With less than 1 wt %, it is found that useful densificationcan not be achieved. While useful densification can be achieved withfrom 1 to 2.5 wt % Al₂ O₃, this can necessitate recourse to closercontrol over heating conditions, as detailed hereinafter with referenceto a mid-temperature hold and/or use of a powder bed which generates asuitable Al species in the sintering atmosphere. It is principally forthese latter reasons that the lower limit of 2.5 wt % is preferred.

Where all of the Mg source is provided in the compact, its level ofaddition (calculated as MgO) may range from about 0.3 to 4 wt %. Thepreferred range (as MgO) is from about 0.5 to 2 wt %, such as from about0.5 to 1.5 wt %. With all of the Mg source provided in the compact asformed, a level of addition (as MgO) less than about 0.3 wt %, the Mgsource appears not to result in generation of a sufficient level ofliquid phase at a low temperature to achieve efficient liquid phasedensification, and an inferior level of densification, comparable tothat achieved with use of Al₂ O₃ alone, results. Above 4 wt % (as MgO),the maintenance of such level of Mg in the final sintered product isdifficult and, if retained, the Mg source tends to result in too high alevel of secondary oxide constituent in the sintered product, withadverse consequences for physical properties. The level of Mg in thesecondary oxide constituent and, indeed, the amount of secondaryconstituent, can be reduced by a longer holding time at the sinteringtemperature, causing such constituent to be lost by decomposition.However longer holding times have adverse consequences for cost ofproduction and tend also to result in undue loss of SiC.

From the foregoing, it will be appreciated that the starting compositionin terms of the Mg source content is related to the Al source content.Alternatively, the overall starting compositions may be expressed interms of the said Al source content. Thus the Mg to Al ratio is in therange 1:2 to 1:25 at a level of 1 to 5 wt % Al source (as Al₂ O₃) andfrom 1:5 to 1:100 at a level of 25 wt % Al source (as Al₂ O₃) Linesconnecting these points define boundaries inside which of the region ofuseful starting compositions are found.

The above indicated levels of Al source, Mg source or both in thecompact can be partially reduced, or fully reduced in the case of Mgsource, by providing that source in a bed or coating to generateAl-species, Mg-species or both in the atmosphere. To the extent thatthis is done, it is difficult to quantify the amount of Al source, Mgsource or both to be provided in the bed or coating. The minimumquantity of Al source, Mg source or both necessary in the bed orcoating, and resultant permeation of the compact by the Mg-species, canvary with both sintering temperature and the rate of heating to thattemperature. Other variables are the thickness of the compact, thevolume of the bed or coating, whether or not a mid-temperature hold isemployed, the volume of the furnace in which the compact is sintered,and the available surface area of the furnace at which condensation ofthe species can occur. However, the quantity of Al source, Mg source oreach required in the bed or coating can substantially exceed thatotherwise required in the compact, such as by a factor of up to 10 oreven up to 20. Also, the quantity can be determined in optimising use ofeach source in a bed or coating by routine analysis of a resultantsample product in a given furnace, and comparison of this with a sampleproduct produced under comparable sintering conditions with each sourcein its compact. As will be appreciated, each source provided in a bed orcoating is required to generate respective species in the furnaceatmosphere at a sufficient partial pressure such that there ispermeation of the compact by the species. If the compact as formed doesnot have at least 0.3 wt %, preferably at least 0.5 wt %, of MgO and therequired level of Al₂ O₃ the level of MgO and/or Al₂ O₃ is to beattained in the compact by permeation. However, the level of MgO inparticular, but also that of Al₂ O₃, can decrease with holding at thesintering temperature.

Instead of use of a powder bed or coating, an external supply ofatmosphere to the furnace can be used to provide Al-species, Mg-speciesor both in the furnace atmosphere. A further alternative is to maintaina favourable ratio of mass of the compact or compacts to the capacity ofthe furnace in which the firing is conducted, such that significantinteraction occurs between individual compacts and between the compactsand the atmosphere. The amount of compacts present relative to thatcapacity is to be such that loss of the densification aids is limited toan extent such that the densification process is completed in areasonable time, avoiding resultant products containing a high level ofresidual porosity. That is, to the extent that the densification aidsare lost, they provide a suitable atmosphere in which densificationoccurs, by generating a sufficient vapour pressure of decompositionproducts of the densification aids. The relative loss of thedensification aids from the compacts (and hence the amount of those aidsretained in the compacts compared to that required to stabilise theatmosphere) is thereby reduced by the mass of the compacts and hence ofthe densification aids initially present. It is preferred that the ratioof the effective furnace volume to the volume of compacts therein be notmore than 5 to 1. However, depending on other factors, it can be as highas 10 to 1 or even as high as 20 to 1. The optimum ratio will bedependent on several factors, including the ratio of compact mass tocompact surface area, heating rate and overall furnace design.

In general, the balance of the compact after allowance for the Al sourceand, if provided therein, the Mg source, substantially comprises SiC.Commercial SiC, which typically is used, has a surface layer of SiO₂ ofup to about 2 wt %. In contrast to some prior art processes fordensification of SiC, it is not necessary to reduce or substantiallyremove such level of SiO₂. Indeed, indications are that the SiO₂ isbeneficial in that, with Al₂ O₃ and MgO, it is necessary to achieveformation of a stable transient liquid at low temperatures. With asource of SiC with an insufficient level of SiO₂, it can be necessary toinclude a small quantity of SiO₂ to the powder mix from which thecompact is formed. Also, at least with higher levels of Al₂ O₃ aboveabout 10 wt %, there can be benefit in adding SiO₂ so as to achieve alevel thereof in excess of 2 wt %, such as up to about 4 wt %. However,a higher level of SiO₂ presupposes that a small proportion of glassyphase can be tolerated in the sintered product, or that the sinteringconditions will be such that a glassy phase is obviated by decompositionof the SiO₂.

The product of the invention may contain Mg, such as at a level inexcess of 0.01 wt %. However, while the Mg level typically is minor, itcan be at a level in excess of 0.1 wt % or higher, such as at a level atleast 1.5 wt % to about 3.5 wt %. Where Mg is present at a level inexcess of an Mg to Al ratio greater than about 1:8, it generally ispossible to establish that the Mg is present as an oxide, in combinationwith Al oxide as an Al rich secondary oxide constituent. Alternativelythe product of the invention may be defined in terms of Mg to Al ratios.Thus the Mg to Al ratio is up to 1:3 at a level of 5 wt % Al source (asAl₂ O₃) and up to 1:6 at a level of 30 wt % Al source (as Al₂ O₃) Linesconnecting these points define boundaries of final composition insidewhich of the region of useful products are found.

The product of the invention preferably is substantially free ofelemental Si. It also preferably is substantially free of a glassyphase. The process of the invention preferably is conducted so as toresult in a product in accordance with one, most preferably both, ofthese requirements. With respect to elemental Si and glassy phase, theupper limit thereof preferably does not exceed about 2 wt % of each.

While preferably substantially free of elemental Si and a glassy phase,the product of the invention may have a secondary oxide constituent asdetailed above. Indeed, the presence of such constituent is preferred,particularly where a product having enhanced fracture toughness isrequired. Sintered SiC produced by prior art pressureless sinteringprocesses typically has a fracture toughness of from 2.5 to 4.0MPa.m⁰.5, while Si-infiltrated or hot-pressed SiC can have a fracturetoughness of from 4 to 5 MPa.m0.5.The present invention enablesproduction of a sintered product, that is, one formed by pressurelesssintering rather than by Si-infiltration or hot pressing, which has afracture toughness in excess of the upper limit of 4 MPa.m⁰.5 forconventionally sintered SiC. That is, the present invention enablesattainment of a fracture toughness level comparable to, or higher, thanthat of Si-infiltrated or hot pressed SiC, and significantly better thansolid state sintered SiC. A product according to the inventionpreferably has a fracture toughness in excess of 4, most preferably inexcess of 4.5 MPa.m⁰.5, such as in excess of 5.0 MPa.m⁰.5.

The fracture toughness of a product according to the invention increaseswith the level of secondary oxide constituent present. In general, atleast about 4 to 5 wt % of oxide constituent is necessary in order toachieve a fracture toughness in excess of 4 MPa.m⁰.5. Above that levelof oxide constituent, fracture toughness can increase to about 4.5MPa.m⁰.5 and 5.5 MPa.m⁰.5 at oxide constituent levels of about 8 wt %and 15 wt %, respectively. An advantage of the present invention is thatit enables the attainment of such level of oxide constituent, andresultant enhanced fracture toughness, for reasons detailed above and inthe following.

The fracture toughness values indicated above for the present inventionwere determined by indentation, using the equation of Antsis et al, J.Amer. Ceram. Soc. 64 [9] 533-538 (1981), using a Vickers Hardnessdiamond indentor and a load of 49N. However, as indicated herein withreference to some Examples, determinations with a standard load of306.6N, using the equations of Niihara et al, J. Mater Sci. Letters 1(1982) 13-16, for Palmquist and median cracks, give numerically highervalues ranging up to at least 5.6 MPa.m^(1/2).

In the above detailed explanation of the mechanism thought to beinvolved in the process of the present invention, the indicatedinvolvement of MgO is explained. Central to this is the formation of anMgO--SiO₂ --Al₂ O₃ liquid phase at a low temperature, with that liquidbeing increased in volume by dissolution of further Al₂ O₃ and beingretained to higher temperatures. Despite this, and despite also that theMgO can be formed in situ from Mg species generated in or supplied tothe atmosphere, it is found that MgO progressively decomposes and islost to the sintering atmosphere with still higher temperatures, andtime at temperature. While the loss of MgO by this means can besubstantially complete, resulting in the low levels of Mg in the productof the invention, the MgO is found to have fully served its function inachieving good densification prior to its loss.

In addition to loss of MgO by its decomposition, it is found that Al₂ O₃similarly can be lost by decomposition, as also is the case in the priorart where Al₂ O₃ is used alone. However, due to the transient liquidphase resulting from the MgO, the loss of Al₂ O₃ is substantially lessthan occurs with use of Al₂ O₃ alone. That is, the activity of Al insolution in the liquid phase is reduced such that decomposition of Al₂O₃ tends to occur principally with that portion thereof, if any, whichis not taken into solution in the liquid phase. The proportion of Al₂O₃, if any, which is not taken into solution is substantially less withuse of MgO compared with use of Al₂ O₃ alone.

Due to any loss of Al₂ O₃ and also the loss of MgO, the product of theinvention can exhibit a weight loss relative to the weight of thecompact from which it is formed. Some weight loss also can result fromdecomposition of SiC and, as the sintered product typically exhibits nodetectable elemental Si or glassy phase, it appears that weight lossalso occurs by decomposition of SiO₂. It appears that the bed orcoating, if containing SiC, also serves the purpose of supplyingvolatile species such as SiO which inhibit the decomposition of SiC inthe powder compact. The loss of Al₂ O₃ can be regulated by generatingAl-species in the atmosphere in which the compact is sintered, or byproviding an atmosphere containing such species, and this if found to bebeneficial. Thus, in one preferred form of the invention, the compact isheated to the sintering temperature in the presence of a bed or coatingof an Al-containing material which generates vapour of the Al-speciessimilar to that formed by the decomposition of the Al₂ O₃ of thecompact. On heating of the bed or coating, the Al₂ O₃ provided or formedtherein is decomposed, with its decomposition products providing therequired Al-species. The quantity of Al₂ O₃ decomposed from the bed orcoating is such that the Al-species are present at a high partialpressure in the furnace and act to prevent decomposition of the Al₂ O₃content of the compact. The principal relevant Al-species is thought tobe Al₂ O.

It appears that generation of Al-species in the atmosphere duringsintering regulates the decomposition of Al₂ O₃ in the compact, at leastto the extent of reducing the rate of that decomposition. However, it ispossibly only to speculate on this. This is because, although a weightloss from the compact occurs at a higher level in the absence ofAl-species being generated in the atmosphere, generation of Al-speciesin the atmosphere can result in the sintered product having an increasedweight percent of Al relative to the SiC content, than was present inthe compact from which the product was formed. Thus, there can be aweight gain with respect to aluminium, even allowing for the loss if anyof MgO and SiC by sublimation or decomposition. Assuming that therelevant Al-species is Al₂ O, this evidently permeates the compact priorto full densification and evidently is converted therein to Al₂ O₃ byreaction with SiO₂ of the liquid phase.

The mechanism by which Al-species result in a weight gain with respectto Al appears to be quite distinct from that involved in the transportof Mg-species to the compact where the Mg source at least partially isprovided in a bed or coating. That is, it seems clear that theMg-species comprises elemental Mg vapour, whereas it is quite improbablethermodynamically that decomposition of Al₂ O₃ will result solely ingeneration of elemental Al vapour. Generation of Al₂ O appears to besubstantially more likely, as could be expected from thermodynamicconsiderations.

Apart from the requirement for generation of Al-species, and alsoMg-species where the Mg source is not provided solely in the compact,the atmosphere in which the compact is sintered preferably is inert.Other constituents of the atmosphere may comprise nitrogen, argon,helium or carbon monoxide. The atmosphere has a low oxygen partialpressure as, for example, created by a graphite furnace or carbon in apowder bed or coating.

As detailed above, with reference to a favourable ratio of the mass ofthe compacts to the furnace capacity, the compacts themselves provide astable atmosphere conducive to densification. This is accomplished bythe amount of densification aids in the compacts at the sinteringtemperatures being in excess of the limit below which densification willnot occur in a reasonable time. The loss of the densification aids tothe furnace atmosphere to provide a stable environment is to be suchthat the amount remaining in the compacts is sufficient to enable thecompacts to be densified. This is dependent on the charge of compacts inthe furnace chamber and the rate of any effective removal of the activedensification aids from the reaction zone.

The process of the invention enables densification over a temperaturerange of from 1500° to 2300° C. However, over the lower part of thatrange to about 1700° C, substantially complete densification cannecessitate recourse to either application of pressure or an increase intime at temperature, or both. Pressureless sintering is preferred and,under these conditions, the onset of densification for at leastpreferred compositions is about 1700° C. Rapid densification, enabling asintering time of about 0.25 to 3 hours at temperature, commences atabout 1900° C. The preferred range for sintering temperature is fromabout 1900° to 2100° C. Sintering temperatures above about 2100° C. tendto increase the extent of loss of Al₂ O₃, MgO and SiO₂ by decomposition,with the loss of Al₂ O₃ being unable to be offset by Al-species in theatmosphere. However, sintering temperatures of from about 2100° to 2300°C. enable a reduction in the time required at temperature, to less thanabout 1 hour. Also, depending on the properties required in the endproduct, substantially complete loss of Al₂ O₃, MgO and SiO₂ can bebeneficial.

The indicated role of MgO, in forming a transient liquid phase at arelatively low temperature, is confirmed by the benefit found to beobtained by a mid-temperature hold. As stated above, it appears that theliquid phase initially forms at about 1300° to 1400° C., and is stableand retained at temperatures above that range. A hold in that range,preferably at the upper end thereof, or slightly above that range, isfound to enhance the level of densification. It appears that soaking atsuch mid-temperature range allows the liquid formed to equilibrate,possibly by allowing increased dissolution of Al₂ O₃ in the liquid and aresultant increase in the volume of the liquid. A hold of from 20 to 180minutes or more, such as about 60 minutes, typically is sufficient forthe purpose of enhancing densification. A similar enhancement canhowever be achieved by a relatively slow rate of heating to thesintering temperature, such that the compact remains in the temperaturerange of from about 1200° to about 1550° C. for a period of about 30 to120 minutes.

The benefit of a mid-temperature soak in enhancing densificationsimilarly tends to confirm that the MgO acts as a flux, facilitatingdissolution of Al₂ O₃ In this regard, it is found that such soak doesnot provide enhanced densification when Al₂ O₃ is used in the absence ofMgO. Indeed, use of Al₂ O₃ alone is found to achieve a lesser degree ofdensification than use of Al₂ O₃ and MgO in combination, at least in agiven time at a given sintering temperature, even where suchmid-temperature soak is not used with that combination. The indicationsare that, as suggested by the Al₂ O₃ --SiO₂ binary phase diagramconsidered above, densification with Al₂ O₃ alone achieves insufficientliquid necessary for densification by liquid phase sintering, at leastuntil relatively higher sintering temperatures are attained. Inaddition, at those relatively high temperatures, the tendency fordecomposition of SiO₂ is increased, thus limiting the potential volumeof liquid phase that can be formed. Also, longer sintering times at suchhigher temperatures are necessary, possibly due to decomposition of Al₂O₃ competing with liquid formation.

The Al-species, Mg-species or both, whether by means of a bed or coatingin the furnace or from an external supply to the furnace atmosphere,most preferably are generated by heating Al₂ O₃, MgO or both. However,with use of these or other sources, the source of the species preferablyis in particulate form, such as grit or powder. The particulate sourcemost preferably includes particulate SiC, as decomposition of this andgeneration of Si-species in the atmosphere passed to or generated in thefurnace is found to assist in minimising decomposition of SiC in thecompact. Also, particulate C in the particulate source is found to bebeneficial, with the C having two benefits. First, the C acts tominimise the tendency for the particulate source to fuse into a bondedmass by reducing the tendency for the formation of elemental Si. Second,the C assists with decomposition of the constituents of the particulatesource and, hence, generation of Al-, Mg-, and Si-species. As analternative to use of particulate C, the particulate source can beheated in a carbon box or vessel of other form. However, use ofparticulate C is preferred.

Where the particulate source is to generate Al-species, but notMg-species, a suitable source comprises from 10 to 85 wt % SiC, from 1to 90 wt % C, and from 1 to 50 wt % A1 source (calculated as Al₂ O₃).Where the particulate source is to generate Mg-species, a similar sourcebut in which Mg source (calculated as MgO) is substituted for Al sourcecan be used. Where both Al- and Mg-species are to be generated, asuitable particulate source comprises from 10 to 85 wt % SiC, from 1 to90 wt % C, and from 1 to 30 wt % of each of Al and Mg source(respectively calculated as Al₂ O₃ and MgO but not necessarily at thesame level of addition).

Our research suggests that the diffusion of Al- and Mg- species into andfrom the compact can take place through the liquid phase. This acts as arapid diffusion path. As will be appreciated, diffusion can take placeat a greater rate through a liquid or amorphous phase than through thecorresponding crystalline form of the same chemical composition.Microstructural studies have revealed that the Al rich secondary oxideconstituent involved in the procedure of the present invention isinterconnected and can result in a duplex type structure. The observedanomaly where the fired bulk density is high but the residual amount ofMg is low is explained in terms of the microstructure. That is loss ofthe MgO and Al₂ O₃ can proceed through diffusion of these species to thesurface of the compacts and their subsequent loss through vaporisation.This mechanism can operate in a dense body. The only requirement is thatthe level of densification aids is sufficient to allow densification toproceed to completion. Thus the observed behaviour is explained. Itshould also be appreciated that the secondary oxide constituent canexist over a range of Al rich compositions as a solid solution. Theflexibility of the ratio of Mg to Al of this phase means that the lossof Mg can be accommodated. This mechanism could also be applied toreduce the amount of secondary oxide constituent after densification ofthe bodies.

The powders to comprise the compact can be prepared by conventionalprocedures, as can formation of the compact to a required shape. Thepowders can be blended by techniques such as wet or dry ball milling.Wet milling can be carried out in water or in a suitable organic fluid,such as iso-propanol. The resultant slurry, containing a required binderas is conventional, then is dried where this is required, such as byspray drying. The powder mix then is formed or compacted into therequired shape by conventional ceramic forming techniques, such asuniaxial pressing, isostatic pressing, a combination of uniaxial andisostatic pressing, injection moulding or extrusion, with thesetechniques being supplemented, if required, by green machining, slipcasting, pressure slip casting or tape casting.

When a binder is used, the compact in its green state most preferably isheated slowly to a relatively low temperature, either in the sinteringfurnace or as a separate operation. This is desirable to effect binderburn-out for removal before the densification step. Such heatingpreferably is at a temperature of from 300° C. up to 550° C. to 700° C.,for a period of from 30 to 90 minutes or more, such as up to 300minutes. However, the temperature and time can vary, depending on thesize of the compact and also on the chosen binder. The heating can be inan oxygen containing atmosphere, such as air, or in an atmospheresubstantially free of oxygen, such as argon or nitrogen; a choice onthis depending in part on the binder type. For a binder which is burntout or removed by decomposition, without leaving a residue, either typeof atmosphere is acceptable. However some binders, such as those leavinga residue typically of carbon, necessitate use of an oxygen containingatmosphere where, as is preferred, the residue also is to be burnt out.

Where the compact includes an organic precursor of Al source, Mg sourceor both, this will typically result in a residue in the compact afterbinder burn-out. However, the benefit of such precursor, such as inacting as a lubricant during powder compaction, will not be offset bysuch residue. That is, the residue will comprise Al₂ O₃, MgO or both,able to comprise at least part of the required sintering aids. Also,such residue increases the Young's modulus of the compact after theburn-out, and can also increase its strength, allowing for easierhandling and green machining of the compact.

The present invention will be further illustrated by examples in anon-limiting manner. In each of the Examples, dimensions indicatedrelate to the compact (that is, after shaping or pressing, and prior tocold isostatic pressing where applicable, the unfired green body), whiledensity values are those of the fired body. Also, an organic binder isused in all Examples, at a level of 2 wt % unless otherwise indicated.Additionally, analysis of sintered bodies are based on the assumptionsthat determinations of C, Al and Mg are respectively attributable toSiC, Al₂ O₃ and MgO; with resultant minor departures from 100% onaggregate.

EXAMPLES 1 to 6

The raw materials used are shown in Table 1.

                  TABLE 1                                                         ______________________________________                                        Starting Materials                                                            Raw Material                                                                            Source  Grade        Treatments                                     ______________________________________                                        α SiC                                                                             Lonza   UF10/UF15    --                                             α Al.sub.2 O.sub.3                                                                Alcoa   A16SG        --                                             MgO       Ajax    Analytical Grade                                                                           Calcined at 900° C.                     ______________________________________                                    

The powders were weighed, and then ball milled with SiC balls. Theconditions used are shown in Table 2.

                  TABLE 2                                                         ______________________________________                                        Conditions Used for Ball Milling Operation                                    ______________________________________                                        Time              16 hours                                                    Powder            300 g                                                       Balls             1500 g                                                      Fluid             600 ml iso-propanol                                         Binder            2 wt %.                                                     ______________________________________                                    

After milling, the balls were removed and the slurry was subsequentlyspray dried. The resultant powder mix was uniaxially pressed intocylinders 38 mm diameter and 33 mm high and then cold isostaticallypressed at a pressure of about 150 MPa. The samples were then heated inair at 30° C.hr⁻¹ to 400° C. and held for 60 minutes to remove thebinder.

The samples were covered by a powder bed in a graphite work box. Thecomposition of the powder bed was 76 wt % Sic grit; 19 wt % Al₂ O₃powder; and 5 wt % C. The work box was then heated in a graphiteresistance furnace in an argon atmosphere. The firing cycle employed isshown in Table 3.

                  TABLE 3                                                         ______________________________________                                        Firing Cycle                                                                  ______________________________________                                        Ramp at 20° C. min.sup.-1                                              Heat to 1400° C.                                                       Dwell for 60 minutes                                                          Ramp at 5° C. min.sup.-1                                               Heat to 2030° C.                                                       Dwell for 60 min.                                                             Cool at 10° C. min.sup.-1                                              ______________________________________                                    

The results obtained are shown in Table 4. In addition to the detailsprovided therein analysis by XRD techniques revealed an increasingamount of Al rich secondary oxide constituent with increasing initialaluminum content, while no α-Al₂ O₃ (corundum) was detected. The SiCcontent in the final body was ascertained, by combustion analysisdetermination of the C content of the body, assuming all the C wasassociated with the SiC. The Al and Mg analyses were performed usingAtomic Absorption techniques on the fired bodies. The elemental analysiswas converted to the equivalent amount of oxide. The Vickers HardnessNumbers were determined employing a load of 5 kg force.

It can be seen that at higher levels of aluminium and magnesiumadditions in the molar ratio of 2 to 1 good properties were not obtainedfor the thick bodies of these Examples; compare Example 2 with Example4. Moreover, the properties were not improved with such ratios havinghigher levels of MgO. The body of Example 4 was cracked on removal fromthe furnace. At still higher levels of additions of magnesia (Example6), the body was cracked and its interior had an inferior hardnessnumber. It is postulated that at high additions of magnesia, itssubstantial depletion is required to achieve excellent properties.Furthermore, it is speculated that the decrease in properties especiallyin the interior of thicker bodies could be of a result of the inabilityof the thicker bodies to substantially deplete the amount of magnesia inorder to achieve a favourable aluminium to magnesium ratio in the body.By contrast it is possible to produce bodies with high aluminium tomagnesium ratios and maintain excellent physical properties with nocracks being detected after firing (compare Examples 2 and 3). It isspeculated that the cracking is caused by either differential sinteringof the surface and the interior or differences in the thermal expansionbehaviour between the inside and the outside of the bodies as a resultof changes in composition during the sintering process caused by anunfavourable Mg to Al ratio in the starting compact.

                                      TABLE 4                                     __________________________________________________________________________    Experimental Results                                                                          Example                                                       Composition     1   2   3   4   5   6                                         __________________________________________________________________________    SiC Surface Area m.sup.2 · g.sup.-1                                                  10  10  10  10  10  10                                        Initial Al.sub.2 O.sub.3 wt %                                                                 3.6 10.7                                                                              23.6                                                                              10.7                                                                              20.7                                                                              17.9                                      MgO wt %        1.4 1.4 1.4 4.3 4.3 7.1                                       Final SiC wt %  94.1                                                                              87.8                                                                              73.1                                                                              73.1                                                                              72.8                                                                              85.8                                      Al.sub.2 O.sub.3 wt %                                                                         4.6 12.6                                                                              24.0                                                                              20.2                                                                              24.9                                                                              13.4                                      MgO wt %        0.5 1.2 1.4 3.9 4.0 2.5                                       Density g · cc.sup.-1                                                                3.10                                                                              3.23                                                                              3.25                                                                              3.24                                                                              3.22                                                                              3.20                                      Macrocracking   No  No  No  Yes Yes Yes                                       Weight change % -2.8                                                                          -3.1                                                                              -1.5                                                                              -4.7                                                                              -2.4                                                                              -4.3                                          Vickers Hardness (GPa) - Edge                                                                 24.3                                                                              27.4                                                                              23.9                                                                              25.9                                                                              20.9                                                                              25.7                                      Centre          13.6                                                                              24.3                                                                              22.2                                                                              20.5                                                                              21.1                                                                              14.1                                      __________________________________________________________________________

EXAMPLES 7, 8 and 8A

Using the procedure of Examples 1 to 6, except as indicated for Example7 and 8A, discs 75 mm in diameter and respectively 7 mm and 13 mm thickwere prepared. The results are shown in Table 5.

                  TABLE 5                                                         ______________________________________                                        Experimental Results                                                          for 75 mm discs 7 mm thick                                                    Example           7        8        8A                                        ______________________________________                                        Composition                                                                   SiC Surface Area m.sup.2 · g.sup.-1                                                    10       10       10                                        Initial                                                                       Al.sub.2 O.sub.3 wt %                                                                           3.6      10.7     10.7                                      MgO wt %          1.4      1.4      1.4                                       Final                                                                         SiC wt %          92.8     86.1     84.8                                      Al.sub.2 O.sub.3 wt %                                                                           8.1      15.6     13.3                                      MgO wt %          0.01     0.01     0.5                                       Density g · cc.sup.-1                                                                  3.19     3.21     3.24                                      Weight change     -0.7     -1.1     -0.8                                      Vickers Hardness (GPa)                                                        Edge              22.6     24.7                                               Centre            23.8     24.2                                               Fracture Toughness (MPa · m.sup.0.5)                                                   4.4      5.4                                                ______________________________________                                         Notes:                                                                        (1) For Example 7, heating from 1400° C. to the soak temperature       was at 20° C. min.sup.-1.                                              (2) Fracture toughness was determined by indentation using the equation o     Anstis, Chantikul, Lawn and Marshall, J. Amer. Ceram. Soc. 64 [9] 533-538     (1981), using a Vickers Hardness diamond indentor at a load of 49N.      

For Example 7, only SiC and a small amount of Si was detected by XRD.For Example 8, a small amount of α-Al₂ O₃ (corundum) was detected.

Table 5 illustrates the high level of fracture toughness achievable in afired body of a sintered SiC product of the invention, and the generaltendency for fracture toughness to increase with the level of Al richsecondary oxide constituent. In contrast, commercial HEXOLOY (SiCdensified by use of B or a B compound) was found to have a fracturetoughness of about 3.1 MPa.m⁰.5 as determined by the Vickers indentionmethod used for Examples 7 and 8, which accords with published data onHEXOLOY.

BRIEF DESCRIPTION OF THE DRAWINGS

Photomicrographs, representative of the microstructure of the sinteredbody of Example 8A are shown in FIGS. 1 to 3, in which:

FIG. 1 is a photomicrograph X 2500 of a polished, unetched section;

FIG. 2 is a photomicrograph X 2000 of a polished and etched section; and

FIG. 3 is similar to FIG. 2 but at X 5000.

FIG. 1 illustrates the form of duplex microstructure obtainable with thepresent invention. The light coloured constituent comprises Al richsecondary oxide also containing Mg, forming a skeletal structurethroughout the sintered SiC. That secondary oxide constituent is notapparent in FIGS. 2 and 3, due to etching, although the equi-axed,rounded grains of sintered SiC are readily apparent. Also apparent isthe absence of very fine grains corresponding to finer particle sizes ofthe SiC powder used, indicative of these having been dissolved in theliquid phase during sintering and precipitation on and between largerparticles. Such precipitation also is apparent from the rounded aspectof the SiC grains, and the absence of sharp edges.

EXAMPLES 9 and 10

In these samples, Al₂ O₃ (AKPHP) from Sumitomo was the sole sinteringaid. The samples were covered by a powder bed in a graphite work box.The powder bed was a blend of 80 wt % SiC grit and 20 wt % Al₂ O₃powder. The work box was heated in a graphite resistance furnace in anargon atmosphere. The firing cycle employed is shown in Table 6, and theresults are shown in Table 7.

                  TABLE 6                                                         ______________________________________                                        Firing Cycle                                                                  ______________________________________                                        Ramp at 20° C. min.sup.-1                                              Heat to 1400° C.                                                       Ramp at 5° C. min.sup.-1                                               Heat to 2030° C.                                                       Dwell for 60 min.                                                             Cool at 10° C. min.sup.-1                                              ______________________________________                                    

                  TABLE 7                                                         ______________________________________                                        Experimental Results                                                          (Al.sub.2 O.sub.3 Only)                                                       Example             9       10                                                ______________________________________                                        Composition                                                                   SiC Surface Area m.sup.2 · g.sup.-1                                                      10      10                                                Sample Dimensions                                                             Diameter mm         75      38                                                Thickness mm        7       33                                                Al.sub.2 O.sub.3 wt %                                                         initial             3.6     10.7                                              final               4.0     11.6                                              Density g · cc.sup.-1                                                                    3.01    2.99                                              Weight change %     -0.9    -1.3                                              Vickers Hardness (GPa)                                                        Edge                18.5    24.4                                              Centre              17.3    12.3                                              ______________________________________                                    

Examples 9 and 10 exhibit an inferior density compared with thatobtainable under the same conditions with use of MgO. Also, comprison ofExample 10 with Example 9 illustrates a tendency for decreasinguniformity with section thickness and Al₂ O₃ level.

EXAMPLES 11 to 13

Samples were prepared under different conditions in regards to thepowder bed and firing cycle for discs 7 mm thick. Details and resultsare listed in Table 8.

                  TABLE 8                                                         ______________________________________                                        Experimental Results                                                          Example         11         12      13                                         ______________________________________                                        Composition                                                                   SiC Surface Area m.sup.2 · g.sup.-1                                                  15         15      15                                         Initial                                                                       Al.sub.2 O.sub.3 wt %                                                                         3.6        3.6     3.6                                        MgO wt %        1.4        1.4     1.4                                        Composition of Powder Bed                                                     SiC wt %        80         73      76                                         Al.sub.2 O.sub.3 wt %                                                                         20         18      19                                         C wt %          0          9       5                                          Firing Cycle                                                                  Hold at 1400° C.                                                                       no         no      yes                                        Density g · cc.sup.-1                                                                3.10       3.15    3.21                                       Weight change   -4.3       -7.1    -1.6                                       ______________________________________                                    

As shown by Table 8, densification is enhanced by use of C in the bed,and also by use of a mid-temperature hold.

EXAMPLES 14 AND 15

Samples were prepared without MgO in the samples, but with MgO added tothe powder bed. The powder bed (composition "A") comprised 75.5 wt %SiC, 19 wt % Al₂ O₃, 5 wt % C and 0.5 wt % MgO. The conditions were asfor Examples 9 and 10, and results are listed in Table 9.

                  TABLE 9                                                         ______________________________________                                        Experimental Results                                                          (External MgO)                                                                Example             14      15                                                ______________________________________                                        Composition                                                                   SiC Surface Area m.sup.2 · g.sup.-1                                                      10      10                                                Initial - Al.sub.2 O.sub.3 wt %                                                                   3.6     10.7                                              MgO wt %            0.0     0.0                                               Final - Al.sub.2 O.sub.3                                                                          8.5     12.2                                              MgO                 0.01    0.12                                              Sample Dimensions                                                             Diameter mm         75      38                                                Thickness mm        7       33                                                Composition of Powder Bed                                                                         A       A                                                 Density g · cc.sup.-1                                                                    3.19    3.24                                              Weight change %     +2.4    -1.2                                              Vickers Hardness (GPa)                                                        Edge                25.1    25.7                                              Centre              25.0    25.8                                              ______________________________________                                    

The use of MgO in the powder bed, when no MgO is initially present inthe body, is shown to be useful in producing dense bodies by Examples 14and 15. From the results it can be seen that the use of MgO in thepowder bed has a beneficial effect on the densification and results inthe formation of a uniform body of high density. This is in contrast toExamples 9 and 10, when no MgO was present in the reaction zone. Inaddition, it can be seen that the MgO is effectively retained in thethicker body (Example 14) after the firing process but is lost from thethinner body (Example 15). It appears that this loss occurs after thedensification is activated. In any case, it is demonstrated that the useof MgO in the reaction zone allows the fabrication of dense uniformthick bodies as evidence by the fired bulk density and the hardnessmeasurements as compared with when it is not present.

EXAMPLES 16 TO 21

Further Examples 16 to 21 were produced by the procedure of Examples 14and 15, at various levels of Al₂ O₃ and with either no MgO or 1.4 wt %MgO in the samples. The powder bed comprised either composition "A", orcomposition "B" having 76 wt % SiC, 19 wt % Al₂ O₃, 5 wt % C and no MgO.Other conditions were as in Table 10.

The results of Table 10 make clear that good densification is possibleif there is MgO in the reaction zone. MgO in the fired body can beslight, although (taking into account experience where no MgO ispresent) indications are that MgO taken up by the body initially is at ahigher level. Thus, with no MgO detected in Example 18, the indicationsare that the relatively thin body enables all MgO to be lost after ithas activated densification. Particularly in the case of thicker samplesof Examples 19 to 21, increasing density with increasing level of Al₂ O₃is apparent. As also detailed herein, fracture toughness also increaseswith increasing content of secondary oxide constituent, such as up to atleast 15 wt % of that constituent.

                                      TABLE 10                                    __________________________________________________________________________    Experimental Results for Discs/Cylinders                                      Example       16  17  18  19  20  21                                          __________________________________________________________________________    Composition                                                                   SiC Surface Area m.sup.2 · g.sup.-1                                                10  10  10  10  10  10                                          Initial - Al.sub.2 O.sub.3 wt %                                                             10.7                                                                              10.7                                                                              25.0                                                                              3.6 10.7                                                                              25.0                                        MgO wt %      0.0 1.4 0.0 0.0 1.4 0.0                                         Final - Al.sub.2 O.sub.3 wt %                                                               16.8                                                                              17.4                                                                              30.0                                                                              5.1 14.5                                                                              25.7                                        MgO wt %      0.01                                                                              0.05                                                                              0.00                                                                              0.01                                                                              1.4 0.15                                        Sample Dimensions                                                             Diameter mm   75  75  75  38  38  38                                          Thickness mm  7   7   7   33  33  33                                          Composition of Powder Bed                                                                   A   B   A   A   B   A                                           Density g · cc.sup.-1                                                              3.19                                                                              3.18                                                                              3.15                                                                              3.15                                                                              3.23                                                                              3.29                                        Weight change %                                                                             +2.7                                                                              +2.5                                                                              +6.9                                                                              -0.5                                                                              -0.9                                                                              +0.4                                        __________________________________________________________________________

EXAMPLES 22 TO 25

Green bodies approximately 130mm square, 19mm thick and weighing about0.5 kg were produced from a powder mix having 10.7 wt %, 1.4 wt % MgO, 2wt % binder (as in previous Examples) and the balance comprising SiC.The powder mix was milled as detailed in Table 2, with the resultantslurry then spray dried. The green bodies were prepared by single endeduniaxial pressing at 35 MPa in a 60 tonne press. They then were heatedin a flow of air at 30° C.h⁻¹ to 400° C. and held at that temperaturefor 4 hours before cooling at 200° C. hr⁻¹ to room temperature.

In respective firings, each of five samples, the green bodies weresintered to produce tiles by heating in an atmosphere of argon or carbonmonoxide supplied to the furnace from an external source and passedthrough the furnace. With each atmosphere, one firing was with a powderbed in the furnace, and another was without such bed. The bed used wasthe same as described in Examples 1 to 6. In each case, the firing cyclewas as set out in Table 11.

                  TABLE 11                                                        ______________________________________                                        Summary of Firing Cycles                                                      ______________________________________                                        Heat at 10° C. min. .sup.-1 to 300° C.                          Hold until vacuum <200 microns                                                Backfill with Ar or CO                                                        Heat at 10° C. min..sup.-1 from 300° C. to 1400° C.      Hold at 1400° C. for 60 min.                                           Heat at 5° C. min..sup.-1 from 1400° C. to 1900° C.      Heat at 2.5° C. min..sup.-1 from 1900° C. to 2030°       C.                                                                            Hold at 2030° C. for 60 min.                                           Cool at 10° C. min..sup.-1 until natural cooling takes                 ______________________________________                                        over                                                                      

For the firings, the green bodies were placed on edge, in a parallelarray in a rectangular, graphite work box. The spacing between tiles wasabout 1 cm, with a spacing of 1.5 cm between the tiles and side walls ofthe box. Where the powder bed was used, a thin layer of carbon black wasprovided on the bottom of the box, with about 2 cm of the bed on thatlayer; with the remaining volume of the box filled with the powder bedso as to result in the bodies being covered to a depth of about 2 cm.Approximately 3 kg of powder bed was used in each firing.

The results obtained with Examples 22 to 25 are set out in Table 12.

                  TABLE 12                                                        ______________________________________                                        Firing Results                                                                                Density (g/cc)                                                                            % Wt.                                             Example   Environment Green    Fired  Loss                                    ______________________________________                                        22a       Ar/bed      1.58     3.26   0.6                                     22b       "           1.57     3.21   0.6                                     22c       "           1.56     3.19   0.5                                     22d       "           1.56     3.21   0.6                                     22e       "           1.59     3.23   0.6                                     23a       Ar/no bed   1.54     3.17   6.4                                     23b       "           1.55     3.19   4.0                                     23c       "           1.53     3.18   3.7                                     23d       "           1.54     3.18   3.9                                     23e       "           1.55     3.17   6.5                                     24a       CO/bed      1.55     3.21   1.2                                     24b       "           1.55     3.17   1.2                                     24c       "           1.56     3.15   0.8                                     24d       "           1.56     3.16   1.0                                     24e       "           1.56     3.20   1.5                                     25a       CO/no bed   1.54     3.21   5.2                                     25b       "           1.55     3.21   3.4                                     25c       "           1.56     3.22   3.6                                     25d       "           1.55     3.22   3.5                                     25e       "           1.55     3.23   5.4                                     ______________________________________                                    

Examples 22 to 25 show good densification, obtained with an atmosphereof argon or carbon monoxide, with or without a powder bed (or comparableexternal source of Al-species or Mg-species).

Chemical analyses were performed by taking a vertical, cross-sectionalslice from selected tiles of each firing of Examples 22 to 25, andanalysing top, middle and bottom portions of the slice. The results aredetailed in Table 13.

                  TABLE 13                                                        ______________________________________                                        Chemical Analyses (Wt %)                                                      Example    Region  SiC        Al.sub.2 O.sub.3                                                                    MgO                                       ______________________________________                                        22c        Top     80.6       14.3  1.0                                                  Middle  84.8       10.3  0.7                                                  Bottom  82.9       13.5  0.8                                       23c        Top     86.5       11.0  1.0                                                  Middle  84.5       10.7  1.0                                                  Bottom  85.8       11.0  1.0                                       24c        Top     83.6       13.1  0.6                                                  Middle  86.4       9.2   0.3                                                  Bottom  84.2       11.8  0.5                                       25b        Top     86.4       11.1  0.7                                                  Middle  85.8       11.1  1.0                                                  Bottom  88.1       11.1  0.9                                       ______________________________________                                    

The results of Table 13 shows good overall results are obtainable withan atmosphere of argon or of carbon monoxide. The tiles show increasinglevels of Al and, to a lesser extent, of Mg (taken respectively to beAl₂ O₃ and MgO in the secondary oxide constituent) from the centre ofeach tile toward its top and bottom, due to the use of a powder bed.This composition variation is not found to have adverse consequences forthe physical properties of the tiles. It is attributed to the relativelylarge mass of tiles in each firing to the effective furnace volume, andthe close confinement of samples in the work box but can be minimised byreducing the quantity of powder bed or its composition, the initiallevel of sintering aid or a combination of these factors. While a powderbed can be beneficial, as evidenced by earlier Examples, Table 13 showsthat use of a bed can be obviated.

EXAMPLES 26 TO 29

Further discs, 75 mm diameter and 7 mm thick were produced using eitherα-SiC powder as in the previous Examples, with β-SiC powder (ex-Stark),or with a mixture of these α- and β-SiC powders. Green bodies for thetiles were prepared from a powder mix of 10.7 wt % Al₂ O₃, 1.4 wt % MgO,2 wt % binder, the balance SiC. The powder was milled, spray dried andcompacted as described in Examples 1 to 6, except that the mix forExample 28 was prepared by blending separate α- and β-spray dried powdermixes. The resultant green body compacts then were sintered, using apowder bed and firing cycle as specified in Examples 1 to 6. The resultsobtained are summarised in Table 14.

                  TABLE 14                                                        ______________________________________                                        α- and β-SiC Evaluation                                                        Density g/cc  % Wt                                                Example SiC Type  Green     Fired   Change                                    ______________________________________                                        26      α   1.72      3.23    +0.4                                      27      50α/50β                                                                      1.79      3.21    +0.9                                      28      25α/75β                                                                      1.81      3.07    +0.1                                      29      β    1.86      3.05    -1.2                                      ______________________________________                                    

The green density increased with increasing β content, consistent withthe lower packing density of the α-SiC powder used. The use of 50:50 α-and β-SiC resulted in an excellent fired bulk density. The use of β-Sicalone and the 25:75 mechanical mix of spray dried α-SiC and β-SiCproduced relative lower fired densities. However, the use of β-Sic hasnot been optimised and these lower densities are attributed to this, asopposed to any intrinsic property of β-SiC. No significant differencesbetween use of α- and β-SiC were observed on the basis of chemicalanalyses of the discs, and the lower densities with 25:75 α- plus β-SiCand with β-SiC along are not attributed to differences in finalcomposition of the discs.

EXAMPLES 30 to 34

Discs approximately 75 mm diameter and about 7 mm thick, were producedfrom a powder mix having 10.7 wt % Al₂ O₃, 1.4 wt % MgO, 2 wt % binderand the balance of SiC. The powder mix was milled, dried and compacted,and the resultant green body compacts then sintered as specified forExamples 1 to 6, except as detailed herein. In some instances themilling was in water, rather than isopropanol. Also, in one instance ofmilling in water, the milled slurry was pan dried, rather than spraydried. In some cases, the dried powder was sieved, before compaction.Powder preparation is detailed in Table 15, while firing results areshown in Table 16.

                  TABLE 15                                                        ______________________________________                                        Powder Preparation Processing                                                          Milling       Drying   Secondary                                     Example  Fluid         Process  Treatment                                     ______________________________________                                        30       Water         Spray    Nil                                           31       Water         Spray    Sieved                                        32       Isopropanol   Spray    Nil                                           33       Water         Pan      Sieved                                        34       Isopropanol   Spray    Sieved                                        ______________________________________                                    

                  TABLE 16                                                        ______________________________________                                        Results with TABLE 15 Processing                                                       Thickness     Density    % Wt                                        Example  mm            Fired (g/cc)                                                                             Loss                                        ______________________________________                                        30       25            3.20       2.16                                        31       25            3.19       2.20                                        32        5            3.21       3.44                                        33       25            3.21       3.26                                        34       25            3.20       2.72                                        ______________________________________                                    

The fired bulk density of all discs was high; there being no significantdifferences in Examples 30 to 34 to be attributed to the differentfabrication routes. That is, it was found that milling in water isessentially comparable to that in isopropanol. The relatively highweight losses were attributed to a temperature measurement problem.Milling in water necessitates a higher drying temperature, tending toincrease the formation of agglomerates. Thus, while sieving has not beenfound to be necessary with milling in isopropanol, it is indicated asdesirable after drying of powder milled in water.

EXAMPLES 35 TO 39

In these further Examples, discs comparable to those of Example 8 whereprepared, in general by the procedure of that Example. The extent towhich each of Example 35 to 39 was the same as or different from Example8 is set out in Table 17.

                  TABLE 17                                                        ______________________________________                                        Relativity to Example 8                                                       Example                                                                              Comment                                                                ______________________________________                                        35     The α-SiC used was Lonza UF10, rather than UF15                         as in Example 8.                                                       36     The α-SiC was black Acheson SiC of 10 m.sup.2 /g and                    not less than 97% purity, rather than green                                   Acheson SiC of not less than 98% purity for                                   Example 8                                                              37     As for Example 36, plus the disc in its green                                 state was 105 mm square and 12 mm thick, and                                  formed by uniaxial pressing at 80 MN/m.sup.2.                                 Also, the powder was milled in water with 4 wt %                              binder and agglomerated, with green body binder                               burn-out at 500° C. for 4 hr.                                   38     As for Example 37, but with aqueous milling of                                the powder with 4 wt % binder. Also, the powder                               was blended with 1 wt % magnesium stearate as                                 lubricant to provide after burn-out, additional                               MgO.                                                                   39     As for Example 36, plus the disc was 20 × 40 × 5 mm               and formed by injection moulding, using 18.5 wt %                             binder. Also, powder milling was in water, with                               1 wt % Mg-stearate included as in Example 38.                                 Burn-out was by heating to 600° C. in an inert                         atmosphere to char the binder, followed by                                    heating in air at 550° C. for 1 hr. to complete                        binder removal.                                                        ______________________________________                                    

The densities, hardness and fracture toughness of the resultant tilesproduced by Examples 35 to 39 are summarised in Table 18. The resultsdetailed in Table 18 for fracture toughness (K_(IC)) were byindentation, as determined with a standard load of 306.6N according toNiihara et al referred to above. The 1/a values do not allow associationto particular crack types and the K_(IC) values have been calculated forboth the Palmquist type (in parenthesis) and the median type. The mediantype is more likely for the load used and the fracture toughness valuesobtained.

                  TABLE 18                                                        ______________________________________                                        Physical Properties                                                                    Density as    Hardness K.sub.IC                                      Example  Fired (g/cc)  (GPa)    MPa · m.sup.0.5                      ______________________________________                                        35       3.24          20.1     5.5 (6.6)                                     36       3.21          19.9     5.6 (6.6)                                     37       3.19          19.5     5.3 (6.5)                                     38       3.23          19.9     5.2 (6.7)                                     39       3.19          18.6     5.6 (6.8)                                     ______________________________________                                    

When B and C are used as sintering aids for black SiC of not less than97% purity SiC, the higher Al content of that grade of SiC is found toresult in excessive grain growth. In practice, it is extremely difficultto control the sintering cycle to avoid this. However, the tilesproduced with Examples 36 to 39, each using such commercial SiC, werefound to exhibit a microstructure of uniaxed grains of less than 5 μm.Also, as shown by Table 18, the tiles produced with that material werefound to have similar fired bulk densities, hardness and high fracturetoughness properties as obtained with more expensive green SiC of notless than 98% purity used for the tile of Example 35. These findings,attributed to the quite different sintering mechanism characterising thepresent invention, are of significance given the cheaper and morereadily available less pure grade of SiC.

Examples 35 to 39 also further illustrate the excellent level ofdensification obtainable with the present invention, as well as thelevel of hardness and level of fracture toughness. These Examplesfurther illustrate the ability of the process to accommodate differentmethods of compact preparation, both in respect of powder milling byaqueous and non-aqueous media, and of methods of compaction. Examples 38and 39 also illustrate the ability to utilise organic salts aslubricants, with these in the case of Al- or Mg- fatty acid salts ableto serve a second role, after burn-out, of providing at least part ofthe required level of Al or Mg source of densification aid.

EXAMPLES 40 TO 49

In the above Examples, unless indicated to the contrary, the compactswere prepared with 2 wt % organic binder, with binder removal asdetailed in Examples 1 to 6. Thus, in general, binder removal wasconducted by heating the compacts in a flow of air at 400° C. for 60minutes. The present further Examples provide a comparison between thisand use of an inert atmosphere rather than air, with argon selected asthe inert atmosphere.

Discs 75 mm in diameter and either 7 mm or 13 mm thick were prepared, ineach case with 10.7 wt % Al₂ O₃, 1.4 wt % MgO, organic binder and thebalance of SiC. In Examples 40 to 44 and 49, green Acheson SiC of notless than 98% purity was used with 2 wt % binder, while in Examples 45to 48, black Acheson SiC of not less than 97% purity was used with about7.5 wt % organic binder and about 1 wt % Mg-stearate. For Examples 40 to44 and 49, the powders were milled and spray dried as for Examples 1 to6 while, in the case of Examples 45 to 48, the powders wereagglomerated.

Example 41 was subjected to binder removal in air in accordance with theprocedure for Examples 1 and 6, to provide a reference for comparisonpurposes. The others of Examples 40 to 49 were subjected to binderremoval in argon, as detailed in Table 19.

                  TABLE 19                                                        ______________________________________                                        Binder Removal in Argon                                                       ______________________________________                                        Heat to 160° C. at 30° C. h.sup.-1                              Dwell at 160° C. for 60 min.                                           Heat 160° C. to 250° C. at 30° C. h.sup.-1               Heat 250° C. to 450° C. at 12° C. h.sup.-1               Heat 450° C. to 600° C. at 30° C. h.sup.-1               Dwell at 600° C. for 3 hr.                                             ______________________________________                                    

Cool at 200° C.h⁻¹ until material cooling takes over.

Each of the discs of Examples 40 to 49, after binder removal, weresintered in a heating cycle, with backfeed with argon, as detailed inTable 11. In this, the discs were in two batches of 5 discs, using agraphite work box for each batch and a powder bed as detailed forExamples 22 to 25.

Prior to binder removal, the green bodies with higher binder contentwere relatively weak and required careful handling. Their strength wasimproved after binder removal, while binder removal was satisfactory forall Examples, indicating the ability to use an inert atmosphere in thisoperation. However, a higher weight loss occurred in this operation withExamples 45 to 48 due to their higher binder content. The weight lossfor Examples 45 to 48 was about 8.4% compared with 1.8% for Example 41(binder removal in air) and about 2.3% for Examples with lower bindercontent removed in argon.

Results obtained with the fired discs are set out in Table 20.

                  TABLE 20                                                        ______________________________________                                        Firings with Binder Removal in Ar                                                     Density (g · cc.sup.-1)                                      Example   Green     Fired     % Wt. Change                                    ______________________________________                                        40        1.72      3.18      -0.37                                           41 *1     1.73      3.09      -0.04                                           42        1.70      3.06      -0.72                                           43        1.70      3.10      +2.44                                           44        1.71      3.18      -0.14                                           45*       1.90      3.16      -0.47                                           46*       1.89      3.11      -0.25                                           47*       1.88      3.09      -1.36                                           48*       1.88      3.11      -0.38                                           49 *2     --        3.21      -0.10                                           ______________________________________                                         *higher binder content, all other lower binder content                        *1  burn out in air, all others in argon                                      *2  green disc 13 mm thick, all others 7 mm thick.                       

The disc of Example 43 was broken on removal from the work box and thecause of this anomaly is unclear. However, apart from this, the overallresults were satisfactory, with similar densities being obtainedregardless of differences in binder removal atmosphere and bindercontent. The densities were lower on average than for most otherExamples, and it is speculated that this is due to the quantity ofpowder bed in combination with the work box size and disc configuration.

COMPARATIVE EXAMPLES (FROM LITERATURE)

In Table 21, there is set out selected detail of comparative specificexamples from Suzuki et al (U.S. Pat. No. 4,354,991), as well as resultsdetailed in the above-mentioned article by Omori et al. The detail fromSuzuki et al is from their Examples 1 to 3 based on use of β-Sic andExample 6 based on α-SiC; the sole sintering aid in each case being Al₂O₃. The detail from Omori et al is in respect of use of Al₂ O₃ alone, orin combination with Y₂ O₃, as sintering aid.

                  TABLE 21                                                        ______________________________________                                        Literature Results                                                            ______________________________________                                        Suzuki et al                                                                  Composition                                                                   SiC - area M.sup.2, g.sup.-1                                                               13       13       13     7                                       Al.sub.2 O.sub.3 - initial wt %                                                            25       15       3      13                                      Al.sub.2 O.sub.3 - final wt %                                                              not given                                                        Temperature °C.                                                                     2000     2000     2000   1950                                    Time min.    300      300      300    300                                     Density g · cc.sup.-1                                                             3.06     3.11     3.13   3.13                                    Omori et al                                                                   Composition                                                                   SiC - area   not given                                                        Al.sub.2 O.sub.3 - initial wt %                                                            10            5                                                  Y.sub.2 O.sub.3 - initial wt %                                                             0             5                                                  Al.sub.2 O.sub.3 - final wt %                                                              1.5           1.0                                                Y.sub.2 O.sub.3 - final wt %                                                               0             2.1                                                Temperature °C.                                                                     2100          2100                                               Time min.    30            30                                                 Density - g · cc                                                                  2.4           3.1                                                (estimate)                                                                    ______________________________________                                    

Fired bulk densities obtainable with the present invention are shown inTables 4, 5, 8, 9, 12, 14, 16, 18 and 20. The results indicate thatincreasing the addition of oxides, increases the fired bulk density.This is unexpected and in stark contrast with the findings of Suzuki etal (Table 21) where the opposite trend was observed. In addition, thetimes reported to reach a significantly higher fired bulk density weremuch longer (see Table 21). In the work of Omori et al, the addition ofAl₂ O₃ only was ineffective for the attainment of high fired densitybodies. It appears that the difference in the results of Omori et al andSuzuki et al is due to use of a powder bed in the work of Suzuki et al.Furthermore, the use of MgO as disclosed herein, provides a very stableliquid which is able to retain the active densification aids which inturn greatly enhances the observed fired bulk density of productsespecially when higher levels of Al₂ O₃ are employed.

At lower levels of oxide addition (see Example 1), the samples hadhigher levels of porosity in the centre. This is illustrated by thelower levels of hardness obtained in the centre as compared to the edge.The higher porosity was confirmed by microscopy. A similar observationwas made at the higher levels of magnesium addition (see Example 6).

It is important to note, that for thinner cross sections, it waspossible to densify samples containing lower levels of oxide additive(see Example 7). From the hardness determination, it can be seen thatthere was no significant difference in the hardness for the edge andcentre regions indicating the sample was uniform with respect toporosity. This is in contrast to Example 1.

It is an observed feature that the aluminium content of the body canincrease during densification and this greatly facilitatesdensification. It is anticipated that both aluminium and oxygen aresupplied to the body from the powder bed or an external source.

When high levels of both oxide additives are employed, the additives areessentially retained. This is in contrast to the work of Omori et al,where the densification aids are essentially lost. This is attributed tothe greater stability of the procedure and the densification aids asdisclosed herein.

When alumina only is used as a densification aid, the result is inferiorto the combined use of alumina Al₂ O₃ and MgO (see Examples 9 and 10).The fired bulk densities obtained after 60 minutes at the maximumtemperature were significantly lower than when MgO was also employed.The hardness values of these examples was significantly lower comparedto Examples 7 and 2 respectively. This demonstrates the profound andadvantageous effect of the use of both Al₂ O₃ and MgO on thedensification process.

The effect of the addition of C to the powder bed is demonstrated forExamples 11 and 12. The result was a significant increase in the firedbulk density with the addition of C to the bed. The beneficial effect ofa low temperature dwell is also shown by comparison of Examples 12 and13. It is suggested that the improved density is the result of theformation of stable phases at these temperatures which allow theretention of the sintering aids at high temperatures therebyfacilitating densification.

The use of MgO in the powder bed (or in an externally generatedatmosphere charged to the furnace), when no MgO is initially present inthe body, was shown to be useful in producing bodies (see Example 14).From the results it can be seen that the use of MgO in the powder bedhas a beneficial effect on the densification and results in theformation of a uniform body of high density. This is in contrast toExample 10, when no MgO was present in the reaction zone of thesintering furnace.

Densified bodies can be successfully produced without the use of apowder bed, coating or external atmosphere supply when certain criterionof sample to furnace reaction zone ratio, furnace type, atmosphere andheating rate are satisfied. This has the added advantage of a moresimple process, and improved surface finish for the production ofbodies. The process is essentially unchanged from the case of the powderbed with respect to the finished product.

As is evident from the foregoing, the present invention provides a denseSiC product, and a process for its production, which obviates the needfor use of B, or a B-compound, plus C. The invention is notcharacterised by problems of uncontrolled grain growth associated withsuch prior art proposal, or degradation of physical properties whichresult from residual C in the product. Indeed, exaggerated grain growthis not a feature of the present process. Although there is ampleevidence of solution precipitation on SiC grains occurring in theprocess of the invention, it appears that the solubility of SiC in thetransient liquid phase obtained in the process of the invention is lowand that the interfacial energy balances are favourable. In the processof the invention, very fine SiC particles of the compact are dissolvedin that liquid and thereafter the dissolved SiC precipitates on largerSiC particles to provide a densified product having fine, uniformequi-axed SiC grains which, on average, typically are less than 5 μm insize.

The invention enables liquid phase pressureless sintering of SiC. Itthus overcomes the inherent technical and economic disadvantages of theuse of temperature and pressure for the attainment of dense bodies.Liquid phase sintering promotes much greater mobility of species overmuch larger distances. It can overcome problems in trying to obtain ahomogeneous dispersion of additives in a powder mixing step and canassist in the elimination of defects especially those greater than thecharacteristic grain size of the starting powders.

Also, the presence of SiO₂ as a surface layer on the SiC does notpresent a problem with the invention. Thus, it is not necessary topre-treat the SiC, such as by washing in HF, or to add C to achieveremoval of the SiO₂ by reaction. In this regard, it is worth noting thatSiC is thermodynamically unstable in relation to SiO₂. That is onexposure to air (and especially in the presence of moisture) SiC isoxidised to SiO₂. The oxidation mechanism is a passive one in that aprotective oxide layer is formed. As a consequence of this is that SiCin a highly divided form, such as a powder, can contain an appreciablelevel of oxygen on the surface. Typically levels of oxygen are 0.6 wt %for a powder with a surface area of 10 m².g⁻¹. Powders with smallerparticle sizes have an increased oxygen content, and a practicable limitto the use of extremely fine powders can result from SiO₂ contamination.For the conventional pressureless sintering route there is a limit tothe amount of SiO₂ that can be tolerated. It is generally believed thatthe use of C in these materials is required to remove the surface silicafrom the silicon carbide powder. According to a U.S. Pat. No. 4,123,286to Coppola et al, the maximum SiO₂ content is 2.0 wt % (equivalent to1.1 wt % oxygen) with a preferred maximum of 0.5 wt % SiO₂ (equivalentto 0.3 wt % oxygen). This level is below typical levels found oncommercial SiC powders. While the oxide layer can be removed or reducedin amount, this introduces another processing step and the powder can be"re contaminated" with respect to oxygen by exposure to the atmosphereor during the process to produce dense fabricated bodies.

Moreover, while the use of oxides to densify SiC by pressurelesssintering techniques has previously been disclosed, their use has beenassociated with several disadvantages. Thus the presence of Al₄ O₄ C andAl₂ OC in the final product, either by their deliberate addition orformation, results from the use of CaO and Al₂ O₃ additives during thedensification, but such oxycarbides are highly reactive especially withwater and moisture and should be avoided. Also, compared to the use ofAl₂ O₃ (or Al₂ O₃ precursor) alone, the present invention provides forshorter reaction times and higher fired bulk density bodies, with a muchhigher production rate for a given furnace volume and lower energycosts. In addition, with use of Al₂ O₃ (or its precursors) alone as thedensification aid, increasing the amount of the Al₂ O₃ results in adecrease in the reported fired bulk density, whereas in the presentinvention there is no such observed decrease. This means that it ispossible to substitute some of the expensive SiC powder with lower costoxide additives thereby reducing the cost of producing product withoutimpairing the fired bulk density as well as increasing fracturetoughness by incorporating an oxide secondary phase. The use of β-SiCpowders, as is preferred with use of Al₂ O₃ alone, is avoided by thepresent invention. Also the present invention does not necessitate theuse of SiC of highest purity levels, since commercial SiC of not lessthan 95% purity can be used.

The liquid phase sintering of SiC enabled by the present invention isbelieved to be such that other non-oxide and refractory oxideconstituents can be incorporated. Thus, it is speculated that forexample B₄ C or carbides or borides of the transition metals such asTiC, TiB₂ or TaC can be incorporated in the form of particles, plateletsor whiskers. Also, it is speculated that SiC as particles larger than 10μm, such as resulting from a bimodal grain size distribution, or asplatelets or whiskers, can be incorporated. Similarly, it is speculatedthat refractory oxides, such as partially or fully stabilised, orunstabilixed, ZrO₂ can be incorporated. In each case, incorporation ofsuch constituents in a matrix of sintered SiC according to the inventionis envisaged.

Finally, it is to be understood that various alterations, modificationsand/or additions may be introduced into the constructions andarrangements of parts previously described without departing from thespirit or ambit of the invention.

I claim:
 1. A process for the liquid phase sintering of silicon carbide,comprising the steps of:forming a shaped, consolidated powder bodywhich, not allowing for binder, comprises a powder mixture containing atleast 75 wt % silicon carbide and from 1 to 25 wt % (calculated as Al₂O₃) of a powder comprising a source of aluminum selected from alumina,precursors for alumina, and mixtures thereof, particles of the siliconcarbide powder of said mixture having a surface layer of silica; andheating the body to a sintering temperature of from 1500° C. to 2300°C., in an atmosphere which is substantially non-oxidizing at saidsintering temperature, to form in said heating step a liquid phase and aresultant liquid phase sintered body; the body, in said heating step,being heated in the presence of a source of magnesium which is distinctfrom the source of aluminum and comprises at least one of magnesia,precursors for magnesia, magnesium vapor and combinations thereof,whereby said liquid phase achieves a transient ternary composition inwhich it contains silica, magnesia and alumina and produces secondaryoxide constituent, said liquid phase being such that the sintered bodyis essentially free of oxycarbide.
 2. A process according to claim 1,wherein said silicon carbide powder is substantially free of particleslarger than 10 μm and has an average particle size substantially lessthan 10 μm.
 3. A process according to claim 2, wherein said siliconcarbide powder has an average particle size less than 2 μm.
 4. A processaccording to claim 1, wherein said silicon carbide has a purity of notless than 97%.
 5. A process according to claim 1, wherein said siliconcarbide has a purity of not less than 98%.
 6. A process according toclaim 1, wherein said silicon carbide powder comprises at least one ofα-phase silicon carbide of any polytype, β-phase silicon carbide,amorphous silicon carbide and mixtures thereof.
 7. A process accordingto claim 1, wherein said aluminum source is selected from aluminumoxides, aluminum hydroxide, aluminum monohydrate, aluminum nitrate,aluminum silicates, organo-metallic salts of aluminum comprisingprecursors for aluminum oxide, and mixtures thereof.
 8. A processaccording to claim 1, wherein said magnesium source is selected frommagnesium oxide, magnesium hydroxide, magnesium carbonate, magnesiumnitrate, organo-metallic salts of magnesium comprising precursors formagnesium oxide, and mixtures thereof.
 9. A process according to claim8, wherein said body as formed includes at least portion of a requiredlevel of a magnesium source.
 10. A process according to claim 9, whereinsaid magnesium source comprises an organo-metallic salt of magnesium.11. A process according to claim 10, wherein said salt of magnesium is afatty acid salt which, in addition to providing said source, acts as alubricant in formation of the compact.
 12. A process according to claim11, wherein said fatty acid salt is magnesium stearate.
 13. A processaccording to claim 1, wherein said body contains up to 30 wt % in totalof said sources of aluminum and magnesium, calculated respectively asAl₂ O₃ and MgO, to provide a ratio of Mg to Al source (as Al₂ O₃ in therange from 1:2 to 1:100, with the balance apart from bindersubstantially comprising silicon carbide.
 14. A process according toclaim 13, wherein there is from 1 to 5 wt % Al source (calculated as Al₂O₃) and said ratio is from 1:2 to 1:25.
 15. A process according to claim13, wherein there is from 5 to 25 wt % Al source (calculated as Al₂ O₃)and said ratio is from 1:5 to 1:100.
 16. A process according to claim13, wherein said ratio of Mg to Al is from 1:2 to 1:50.
 17. A processaccording to claim 1, wherein the body contains from 2.5 to 20 wt % ofsaid aluminum source (calculated as Al₂ O₃).
 18. A process according toclaim 1, wherein the body contains from 0.3 to 4 wt % of said magnesiumsource (calculated as MgO), said source being selected from magnesia,precursors for magnesia and mixtures thereof.
 19. A process according toclaim 18, wherein said body contains from 0.5 to 2 wt % of saidmagnesium source (calculated as MgO).
 20. A process according to claim18, wherein said body contains from 0.5 to 1.5 wt % of said magnesiumsource (calculated as MgO).
 21. A process according to claim 1, whereinthe body is heated in Said heating step in the presence of an atmospherecontaining at least one of an aluminum containing vapour species andmagnesium vapour at a vapour pressure whereby the alumina and magnesiacontent, respectively, of the oxide phase is controlled.
 22. A processaccording to claim 21, wherein said body as formed is substantially freeof said magnesium source, and said atmosphere contains magnesium vapourat a vapour pressure sufficient to cause said vapour to permeate throughthe body and thereby form said oxide phase.
 23. A process according toclaim 21, wherein said body is heated in said heating step in thepresence of a particulate bed containing at least one of said aluminumsource and said magnesium source, whereby the respective saidaluminum-containing vapour species and said magnesium vapour isgenerated in said atmosphere during said heating step.
 24. A processaccording to claim 23, wherein said bed contains both an aluminum sourceand a magnesium source provided by particulate material comprising ablend of alumina and magnesia.
 25. A process according to claim 23,wherein said bed contains particulate silicon carbide.
 26. A processaccording to claim 23, wherein said bed contains particulate carbon. 27.A process according to claim 21, wherein said at least one ofaluminum-containing vapour species and magnesium vapour is provided byflow thereof, from an external supply, through a furnace in whichheating step is conducted.
 28. A process according to claim 27, whereinsaid external supply is generated by heating outside said furnace aparticulate material containing both an aluminum source and a magnesiumsource provided by particulate material comprising a blend of aluminaand magnesia.
 29. A process according to claim 21, wherein said body insaid heating step is heated in the presence of a coating formed ofparticulate material and containing at least one of said aluminum sourceand said magnesium source, whereby the respective said aluminumcontaining vapour species and said magnesium vapour is generated in saidatmosphere during said heating step.
 30. A process according to claim29, wherein the body is at least partially enclosed in said coating. 31.A process according to claim 29, wherein said coating contains both analuminum source and a magnesium source provided by particulate materialcomprising a blend of alumina and magnesia.
 32. A process according toclaim 29, wherein said coating contains particulate silicon carbide. 33.A process according to claim 29, wherein said coating containsparticulate carbon.
 34. A process according to claim 21, wherein said atleast one of aluminum vapour species and magnesium vapour is generatedby partial loss and decomposition of aluminum source and magnesiumsource, respectively, from the body during heating in said heating stepwherein the ratio of the mass of the body or bodies to the effectivecapacity of a furnace in which the heating step is conducted is suchthat said vapour pressure is generated by said decomposition and is suchthat the extent of Said loss is limited by said ratio and does notimpair formation of said liquid phase and attainment of said resultantliquid phase sintered product.
 35. A process according to claim 1,wherein in heating to said sintering temperature, the body is retainedin the temperature range of 1200° C. to 1550° C. for a period of from 30to 120 minutes whereby formation of said liquid phase is enhanced.
 36. Aprocess according to claim 1, wherein said sintering temperature is inthe range of 1900° C. to 2100° C.
 37. A process according to claim 1,wherein the heating step is conducted at the sintering temperature for aperiod of from 0.25 to 3 hours.
 38. A process according to claim 1,wherein said oxide is retained in the body on completion of said heatingstep.
 39. A process according to claim 38, wherein said oxideconstituent substantially comprises aluminum oxide.
 40. A processaccording to claim 39, wherein said oxide constituent includes magnesiumat a magnesium to aluminum ratio of not more than 1:3.
 41. A processaccording to claim 1, wherein the heating step is conducted at thesintering temperature and for a period of time at said temperature suchthat said oxide constituent is caused to become substantially depletedin magnesium, and the sintered body is substantially free of magnesium.42. A process according to claim 41, wherein said sintering temperatureand said period of time is such that said oxide constituent is caused tobecome substantially depleted in aluminum, and the sintered body issubstantially free of any secondary constituent.
 43. A sintered ceramicproduct comprising a body of silicon carbide produced by the process ofclaim
 1. 44. A sintered ceramic product comprising a body produced bypressureless liquid phase sintering and having at least 65 wt % siliconcarbide, from about 5 wt % to not more than about 30 wt % secondaryoxide constituent, and not more than about 2 wt % each of elementalsilicon and glassy phase; the SiC of said body substantially comprisingequi-axed, rounded SiC grains; said body being essentially free ofoxycarbide and, apart from incidental impurities, free of any rare earthelement, and having a fixed bulk density in excess of 2.95 g.cc⁻¹.
 45. Aproduct according to claim 44, wherein any oxide constituentsubstantially comprises aluminum oxide.
 46. A product according to claim44, wherein any oxide constituent comprises aluminum oxide containingmagnesium at a magnesium to aluminum ratio of not more than 1:3.
 47. Aproduct according to claim 44, wherein said body comprises at least 80wt % silicon carbide.
 48. A product according to claim 47, wherein saidbody comprises at least 85 wt % silicon carbide.
 49. A product accordingto claim 47, wherein said body comprises at least 98 wt % siliconcarbide.
 50. A product according to claim 47, wherein said bodycomprises at least 95 wt % silicon carbide, the body exhibiting amicrostructure in which any constituents other than silicon carbide aresubstantially present in solid solution.
 51. A product according toclaim 44, wherein said body has a fired bulk density in excess of 3.00g.cc⁻¹.
 52. A product according to claim 44, wherein said body has afired bulk density in excess of 3.15 g.cc⁻¹.
 53. A product according toclaim 44, wherein said body has a hardness in excess of 18.5 GPa.
 54. Aproduct according to claim 44, wherein said product has a fracturetoughness in excess of 4 MPa.m⁰.5 (based on the Antsis et al equation).55. A product according to claim 54, wherein said fracture toughness isin excess of 4.5 MPa.m⁰.5.
 56. A product according to claim 44 whereinsaid body has a hardness in excess of 26 GPa.
 57. A product according toclaim 44 wherein said body has a fracture toughness in excess of 5.0MPa.m⁰.5.
 58. A product according to claim 51, wherein said fired bulkdensity is in excess of 3.1 g.cc⁻¹.
 59. A product according to claim 52,wherein said fired bulk density is in excess of 3.25 g.cc⁻¹.